!

_-" NASAContractorReport168218

." _

i. "_ J

¢

_'"_ j

•_-4 LOW-COSTSINGLE-CRYSTALTURBINEBLADES

._ VOLUME1

,;.:_ --

.,': ,.

:._ T.E. Strangman,B. Heath,and M. Fujii i

"; i GarrettTurbineEngineCompany _,

_.--, Phoenix,Arizona85010

- 7.:

{IIkSA-CR-168218) LOW-CQSI ST_GLE-'CRISTAt _84- 15247

_ORB.ISE BL&DESe VOLUBE I _J.nal ReFol:t

(Gannett Tu_b£ne _.ngineCo.) 218 p

_lC _.I0/Ii¥AOI cSCl. II¥ Onclas 6_/26 18043

November1983

Preparedfor

NATIONALAERONAUTICSAND SPACEADMINISTRATION

Lev,,isResearchCenter

UnderContractNAS3-20073
I

!

FOREWORD

_ This Project Completion Report was prepared for the National

Aeronautics and Space Administration, Lewis Research Center. It

presents the results of a program conducted to establish exothermic

heated casting technology for the manufacture of low-cost, single-

;

crystal, uncooled turbine blades for gas turbine engines. The

• program was conducted as part of the Materials for Advanced Turbine

_ _ Engines (MATE) Program under Contract NAS3-20073.

The special contribution of C. M. Phipps, the director of de- i

velopment casting activity at Jet--shapes, Inc. (a Division of Duradyne Technologies), is hereby acknowledued. Mr. Phipps directed :

[ _ the actual foundry activity, provided knowledge and expertise in _

developing this advanced casting process, and contributed to the

_ description of the foundry practices in this report. The authors ,

, - also wish to acknowledge the contribution of G. S. Hoppin Ill, who

_-: originated this project at GTEC and directed its initial phase. _.,_,._

. J

,e

'

1

i_ '_

ti

!

i ,

t
TABLE OF CONTENTS

"%"

SECTION I

.4, 1.0 SUMMARY 1 ,

:_-'- SECTION II

,,,: 2 0 INTRODUCTION 7

_' SECTION III

3 0 SINGLE-CRYSTAL CASTING TECHNOLOGY DEVELOPMENT ii

-'r 3.1 Mold System Modifications 12 "_

;, 3.1.i Crystal Selector Evaluation 12

3.1.2 Wax Assembly Configuration 17

•_ 3 1 3 Ceramic Shell Mold Manufacture 17 _

:__-.. ,j

_, 3.2 gxothermic Heating Improvements 19

-,,,l,:r: 3.3 Casting Parameters Optimization 23 ._

_K-,_, 3 .4 Post-Casting Treatment 25 _

_': SECTION IV

' 4.0 SC MAR-M 247 PROPERTY EVALUATION 27

•,_: 4.1 Tensile Tests 29 ,_

>_ 4.2 Stress-Rupture Tests 33 ._

"2' SECTION V

._; 5.0 SC ALLOY DEVELOPMENT 39

5.1 "Paper Alloy" Study 40 ":_

5.2 Microstructural Stability Study 42 ;_

5.2.1 Alloy 1 45 -_i

5.2.2 Alloy 2 48 _

: _ 5.2.3 Alloy 3 48 :_

5.2.4 Alloy 4 48

5.2.5 Alloy 5 55

PRECEDI:4G rAGE BLANK NOT, FILMEI_

ll..INTENTIONALLY
TABLE OF CONTENTS (Contd)

_ 5.3 SC Derivatives of Mar-M 247 55

5.3.1 Master Alloys 55

5.3.2 SC Casting Campaign- 55

* 5.3.3 Heat Treatment StL'17 59

,_- 5.3.4 Proper ty Characterization 66

5.3.5 Tensile Tests 66

_* 5.3.6 Stress-Rupture Tests 68

_ 5.3.7 Oxidation Evaluation 80

_; 5.3 8 Density 80

5.3.9 Improper Processing 81

_ 5.3 i0 SC Alloy Selection 84

SECTION VI

J

." 6.0 ALLOY PROPERTY CHARACTERIZATION 87 ._-

. 6.1 Test Material Production 87

Z."

*' 6 i.i Master Alloys 87

6.1.2 Blade and Specimen Castings 87

"_ 6.2 Mechanical Properties 95 _ *'_''-

: 6.2.1 Tensile Properties 95

" 6.2.2 Creep-Rupture 102

t 6.2.3 Low-Cycle Fatigue 116

. 6.2.4 High-Cycle Fatigue 125

6.3 Environmental Tests 138

: 6.4 Physical Properties 144

6.4.1 Thermal Expansion 144 ;

& 6.4.2 Thermal Conductivity 145

6.4.3 Elastic Modulus 145

' I

i

SECTION VII i

7.0 8LADE DESIGN 151

7.1 SCope 151

7.2 Preliminary Blade Design 151

_' 7.3 Final Blade Design 152

f

" t

iv
> i

<_ TABLE OF CONTENTS (Contd)

3

No _

, 7.3.1 Mechanical Design 152

.:% 1

,_a 7.3.1.1 Blade Life Analysis 162

_ 7.3.1.2 Blade Vibration Analysis 167

_" 7.3.2 Aerodynamic Design 170

4;-:. 7.3,3 Design Modifications 170 -4a'.

_i" SECTION VIII

_-_ 8.0 COMPONENT MANUFACTURE 173

-'%

_: 8.1 Tooling 173 "

"'_ ? 2 Casting Campaign 173 c

,_: 5.3 Casting Inspection and Testing 176 _

_, 8.4 Blade Machining 177 ,

_ SECTION IX

[ 9.0 COMPONENT TESTING 179 ;

_ 2

_" .

: 9.1 Holographic Testing 179

_- 9.2 Firtree Testing 182 _ _

9.3 Rig Testing 183 _] " -

,_ 9.4 Post-Test Inspection 188 _

SECTION X ,%.

_.: '_

=_ i0.0 COST ANALYSIS 195

; SECTION Xl _

ll.0 CONCLUSIONS 199 ,_

APPENDIX A 202 '_

4

APPENDIX B 203 "_

'._

APPENDIX C 205 _

REFERENCES 215

v L
4

?

:, ORIGINAlP_. ,_'2Ei_ +

OF POOR QUAL.ITY

SECTION I

1.0 SUMMARY

's

J

The demand for more efficient and economical engines continues

to push the industry toward higher turbine operating temperatures.

The ability to keep up with this trend is strongly tied to advance-

_ ments in material technology. Single-crystal (SC) castings offer i

• many significant advantages over turbine materials currently used

: in production engines, such as higher melting point and higher

strength. The goals of MATE Project 3 were to develop a low-cost

casting process capable of producing SC turbine blades and to _ _,_

•!_ demonstrate the capability of the SC blades through extersive •._:

-,_-: property, rig, and engine testing.

'.2 _'_, _ ,

._,_,: The casting process chosen for development was a low--cost, ' L.

_, nonproprietary modification of the exothermic directionally solid- _-...,

" ified (DS) casting process. The similarities between the DS and SC _. ,

7 processes are shown in Figure i. Both processes use exothermic " "

"" DIRECTIONALLY SOLIDIFIED([IS) SINGLECRYSTAL(SC) :

-.

!,

.

' i

STARTER'I/ICRYSTALRIGHT'ANGLEsELECTOR _i

Figure i. Exothermic Casting Process.

%
aluminum/iron oxide nuggets to preheat the ceramic mold above the

melting point of the alloy, and both use a water-cooled copper

chill plate to directionally extract heat from the mold. The major

difference between the DS and SC processes is that the SC process

uses a geometric crystal selector to "select" one of the DS

grains from the starter and then permit this one _roperly

oriented grain to enter and fill the blade cavity. This produces a

casting without grain boundaries.

The initial work focused on characterizing the mechanical

properties of SC Mar-M 247, using SC castings produced with the

Project 1 final design blade mold as a source for test bars.

Testing indicated that properly [001] oriented SC Mar-M 247 did not

have significantly better stress-rupture properties than DS Mar-M

- 247. Therefore, to fully exploit the benefits of single-crystal

castings, a new series of alloys specifically designed for

single-crystal applications was investigated.

Since there are no grain boundaries in a "true" SC casting, "

the grain boundary strengthening elements such as B, Zr, and Hf ce_

be eliminated. The benefits produced by these changes are twofold.

First, by eliminating these elements, which are strong melting point _

depressants, the incipient melting point of the alloy is significantly increased. Second, this higher melting point allows the

alloy to be heat-treated at a higher temperature where all the

eutectic gamma prime can be taken into solution and subsequently

reprecipitated as fine gamma prime. This significantly increases

, the stress-rupture strength of the alloy.

[

Eight single-crystal alloy modifications of Mar-M 247 were

studied analytically, and five of these alloys were evaluated

experimentally through a series of casting trials and screening

tests. Ultimately, two of these alloys were chosen for full

2
property characterization: NASAIR I00 and SC Alloy 3. Figure 2

_ indicates the relative strengths of these two alloys.

2

Using the material property data generated on SC NASAIR 100

and SC Alloy 3, the final blade configuration was designed. This

blade (shown in Figure 3) is essentially the same as the Project 1

_- final design, with the addition of a tip "winglet" and an aft plat-

_ form flow discourager. Analysis of this design indicated that the

_ single-crystal blade has significantly greater LCF and tensile

_ strength than the DS Mar-M 247 blade and more than twice the .

stress-rupture life.

_. Vibrational response of the final blade design was evaluated ._

_ on 12 single-crystal blades using holographic techniques. Results

_', of this testing indicated that the SC blade responded in a similar

_ fashion to the Project 1 DS blade, but at a slightly lower fre-

,[ quency due to the increased mass at the tip of the SC blade. Since :

_ the response of the SC blade was so similar to that of the DS blade,

,* stress coat evaluation of the SC blades was not required. _

• Two complete sets of single-crystal blades were cast and mach- -'

ined to the final blade configuration for component and engine

testing. Twelve of these blades were instrumented with dynamic 'j,

i" strain gauges and assembled into a complete rotor for component

testing (Figure 4) in the high-pressure rotor rig (HRR).

_ This rig allows strain-gauge monitoring of turbine components

operating in an engine environment. The rig test was completed as '

scheduled, but during the routine post-test inspection, four Alloy

3 blades were discovered to have cracks across the leading edge at

about 66-percent span. No other problems were found. Detailed

examination of the cracked blades revealed previously undetected

! high-angle, extraneous grains in the airfoil at the crack location. i

It was determined that the cracks followed these extremely

!

3
' ORIGIDIAL_:_"_''' '+"_

, OF pOOR Q_jAi._T_

LCF

275614001-

+.. _++°c_;Z+O°F_,.,,._7eO°Ci+_'F_

...... +-'-,-..."_--T. _-

+ +3iStZm+"

IIC

- 649oC1_,=, -_ _.,.._

w_

69011001- 104 105

+ CYCLESTOFAILURE

i Fig_tre 2. SC NASAIR 1O0 and SC Alloy 3

Turbine Blade A11Oy Properties.

Ig8400717g-o
OF POOR QUALITY

4.

r _

4

|

S

• " ' _ 'U_

Figure 3. MATE Project 3 SC Blade, Final Configuration.

weak grain boundaries, and they had occurred early in the test. No '0

' signs of fatigue were noted in any of the cracked blades. The

.? remaining blades were carefully reinspected for extraneous grains,

, and several more were found, some of which had run in the HRR test.

All of these blades were eliminated from further testing. After ,,_

reinspection was completed, full-scale engine testing was initiated -,

'l as scheduled. Details of engine testing are presented in Volume 2

of this report. _!

s
Ok,t.i-, _'-<.'_-

OF POOR QUALITY

;' Figure 4. MATE Project 3 Rotor Assembly Showing Locations

of Instrumented Blades.

l

f

5
• SECTION II

2.0 INTRODUCTION

The NASA Materials for Advanced Turbine Engines (MATE) Program

is a cooperative effort with industry to accelerate introduction of

0

new materials into aircraft turbine engines. As part of this ef-

' fort, Garrett Turbine Engine Company (GTEC) was authorized under

NASA Contract NAS3-20073 to develop a new technique for manufac-

• :[

: turing low-cost, single-crystal (SC), uncooled cast turbine blades

" to reduce SC casting costs and improve fuel consumption in advanced i

., turbofan engines. The process development included those efforts

required to transfer the technology from the previously demon-

: _"% strated feasibility stage through component demonstration and ._

"-'; engine test. Portions of the overall effort included p.-ocess

_ scale-up, alloy evaluations, mechanical property generation, hard-

_. ware procurement, component testing, and full-scale engine testing

: to evaluate potential benefits. •....;

f

%

- This report constitutes Volume 1 of a two-volume Project Completion Report presenting the results of the investigations and _

tests performed under MATE Project 3, Low-Cost Single-Crystal Tur-

° bine Blades. This volume covers all of Project 3 except the fullscale engine testing and post-test analysis. These are the

subjects of Volume 2 of this report.

t

The intent of Project 3 was to develop a low-cost process to

;',P produce single-crystal, uncooled turbine blades and to design and

i_ substitute this blade for the solid, directionally solidified (DS)

turbine blade used in the high-pressure turbine of the GTEC TFE731 turbofan engine.

t

7

t-'- _
Project goals associated with this program included the following:

o Development of a nonprop:ietary, low-cost SC casting process

o Improvement of the stress-rupture capability of the SC

alloys re]&_ive _o DS Mar-M 247

o Design of an uncooled HP turbine blade to use the SC material

o Demonstration of uncooled SC turbine blades through component and engine testing.

Project 3 was subdivided into the following seven tasks:

: Task I - Casting Technology

_. Task II - Alloy/Process Selection

Task III - Property Characterization

Task IV - Blade Design

Task V - Component Manufacture and Testing

Task VI - Engine Test

Task VII - Post-Test Analysis

In Task i, the exothermic casting process was modifi¢_ to consistently produce acceptable s&lid SC high-pressure turbine blades

for the TFE731 turbofan engine. This effort focused on tl_e casta_ility and baseline mechanical properties of blades produced with

Mar-M 247.

During Task II, SC alloy derivatives of the Mar-M 247 alloy

were assessed for castability, microstructure, and improvements in

mechanical properties. Two of the3e SC compositions, NASAIR 100

and Alloy 3, were selected for detailed characterization.

!

Ji
In Task III, mechanical, environmental, and physical properties of SC NASAIR 100 and Alloy 3 were quantified. Mechanical pro-

• perty testing included creep rupture, tensile, and high- and low- i

cycle fatigue, and the determination of the effects of protective

coatings. Environmental characterization included coated and

._ uncoated oxidation and hot-corrosion t_sts. Physical properties

measured included density, elastic modulus, thermal expansion, and

[/.

thermal conductivity.

The SC HPT blade was designed in Task IV to utilize the

improved mechanical properties of the SC alloys (NASAIR 100 and

Alloy 3).

__;-

,- During Task V, two full sets of SC turbine blades and the

_'._ required support hardware for both rig testing and the Task VI

I

_.. engine tes.ing were manufactured. Rig and bench testing were also ,_

•_- completed in Task V. _ !_.._

_- ! _-

" The engine testing, test results, and post-test evaluation are •

_ described in Volume 2 of this report. _ ,

: {

!

I

? i

l

I

I

i

I

i
" ORIGINALPAGE |g

OF POOR QUALITY

SECTION III

i .

3.0. SINGLE-CRYSTAL CASTING TECHNOLOGY DEVELOPMENT

A primary program objective was to further develop the non-

: proprietary, low-cost exothermic casting process to produce

_ uncooled single-crystal (SC) high-pressure TFE731 turbine blades of

• ", Mar-M 247 and higher strength derivative alloys. This process was

_- selected based on success achieved in Garrett's NASA-MATE Pro- :

• ject I, "Low-Cost Directionally Solidified (DS) Turbine Blades, "I

and work performed by Jetshapes, Inc. (Rockleigh, New Jersey). : i

This success verified the feasibility of adapting the exothermic

process to preduce SC turbine components. , .z

• ._:

'$L

As illustrated schematically in Figure 5, the DS and SC exo- l

_" thermic processes are conceptually very similar. The casting pro- '

r; .I ....

, DIRECTIONALLY SOLIDIFIED(DS} SINGLECRYSTALISC} 1 .,

ff

' " " '

1

): ?TARTER / RIGHT-iNGLE i !

CRYSTAL 1

, SELECTOR i

,_:;,', i

; Figure 5. Exothermic Casting ?rocess. I

ll <

I!/(_. [O ,,INTENTIONALLBLYM4K

pRF.,CKDINGPAGE BLANK NOT. FILMFJ_

6
!

!

,J

!

cess derives its name from the use of heat released during the excthermic reaction between aluminum and iron oxide finely dispersed

in briquets or nuggets. This heat is used to preheat the ceramic

investment casting shell mold to a temperature above the melting

point of the alloy. After preheating, the mold is placed on a 1

water-cooled copper chill to establish unidirectional heat extracZion conditions. Then, the alloy is cast into the mold and directionally solidified. Figure 5 also shows that the principal modification of the process is the inclusion of a geometric crystal

: selector in the ceramic shell mold at a position between the

starter block and the blade cavity. The crystal selector is

designed to permit only one properly oriented grain to enter the

. blade cavity.

¢

_. Although the selector geometry was an important consideration,

_ process development included additional aspects:

_. o Mold system• modifications

o Exothermic heating improvements

o Casting parameter optimization. "

3.1. Mold System Modifications ?

"p

Development of the mold system included these areas: i

; o Crystal selector evaluation

J o Wax assembly configuration i

o Ceramic shell mold manufacture.

Each aspect of tailoring the investment casting shell mold for SC ii

i., castings is reviewed in the following paragraphs. ! _t

_ : 3.1.1 Crystal Selector Evaluation [

r

l

Based on experience, trials were run with both helical crystal ' I[

I

selectors and selectors using several ramped right-angle turns to

12

.................................. (i)'
Po0 _ '_- _u

QUALITy

choke off excess grains. Jetshapes worked with different helix

diameters and lengths, different DS starter chamber heights, and

different blade orientations to feed metal and crystal entry. The

right-angle starters used one DS originating chamber, feeding grain

into from one to four mold cavities. The geometries of some of the

evaluated crystal selectors are shown in Figure 6.

: The starter block adjacent to the chill is a 25.4 mm (I inch)

high cube. In this starter block, the casting texture transitions

from essentially random to predominantly [001] parallel to the heat i

flow direction. Grains permitted to enter the crystal selector are

taken from the central region of the selector, which best approxi- _

. mates the unidirectional heat flow requirement. Construction of ._

x.

:l

ml n

A

W I

" ._ t

v V'

; %'

E

|,., 50 mm l p6o_oo-9 i" _.

Figure 6. Single-Crystal TFE731 Turbine Blades Cast with RightAng Selectors le Crysta(Rig l htS),elecOne-h torsalf (LeSiz ft)e. and Helix Crystal 1 '

13

t®'
the starter block is consistent with observations by Gell, et al. 2

that grains with a <iii> orientation are quickly overgrown by

grains with <001> and <011> orientations• In that study--with a L

_ similar alloy--it was found that after 20 _m (0•787 in.) of growth,

more than 90 percent of the grains had the [001] direction within

t 15 degrees of the starter block vertical axis.

i

As indicated in Figure 6, the helix selector started from the

• center of the starter block. In contrast, the right-angle selector

;° had an additional 20 mm (0.787 in.) long extension from the starter

block before the grains entered the crystal selector.

o- The function of the c_ystal selector is to select and permit

only one properly oriente_ SC grain to enter the blade cavity of

_ the mold. Grain orientation results for the first 21 molds of SC

<+_+ Mar-M 247 TFE731 HP turbine blades are provided in Table 1 This

;7" table shows that the right-angle selector was more effective than

the helix in producing SC blades. With a right-angle selector, ! +_

• 86 percent of the SC blades produced had the required orientation +

within 12 degrees of the [001] ; with the helix selector, only

t+ 42 percent of the SC blades produced had the desired orientation.

This evaluation also indicated that the best single-crystal yield !

was generated with the blade root toward the chill plate.

r

+" Two variations in helix selector cross-section diameters were

• used in this evaluation: 3.2 mm (0.126 in.) and 4.0 mm (0.157 in.).

Right-angle selectors also included several variations. The principal modification was growing tt_o blades instead of one from each

"'_ selector. This approach minimized the amount of metal consumed in +

the selector. In addition, eight other minor variations in selector geometry were evaluated to arrive at the two-blade, rightangle starter s_lected for the remainder of the program (Figure 7).

(Productivity was expected to be increased eventually in a production process using one starter for four blades.) +

14
_; TABLE i. CRYSTAL GROWTH ORIENTATIONS* OF TFE731-MATE 3

SINGLE-CRYSTAL BLADES PRODUCED IN EXPERIMENTAL

MOLDS 1-21.

i

S ingle-Crystal Blades*

Crystal Within Within Unsatisfactory

Mold Selector Blades 6 ° of 12 ° of Other

No. Configuration Cast [001] [001 ] Or.ientations _

1

: 1 Helix 16 1 1 3

'. 2 Helix 16 0 5 8

3 Right angle 14 4 4 0

- 4 Right angle 16 2 0 0

• 5 Right angle 15 4 0 1

6 Right angle 15 6 3 0

7 Heli x 16 1 1 4

8 Helix 16 2 2 2

9 Right angle 14 2 3 3

i0 Hell x 14 0 0 1

_ ii Right angle i0 5 3 2

12 Right angle I0 4 2 3

[_oj 13 Right angle 12 0 4 1

14 Right angle 12 0 ii 0

" 15 Right angle 12 0 2 0

_ 16 Right angle 12 0 1 0

17 Right angle 12 0 4 A+R (9 A/F) 0 iJ 18 Right angle 12 0 5 A+R (ii A/F) 1 o

/ 19 Right angle 12 0 8 A+R (i0 A/F) 0 ! _

20 Right angle 12 0 5 A+R (7 A/F) 1

21 Right angle 12 0 8 A+R (9 A/F) 1

NOTES: (i) Determinations on Molds 1 through 9, ii, and 12 by i

Laue pattern; on Molds 13 through 21 by visual i

review.

(2) Molds ii and 12 also contained _wo 50.8 mm (2 inch)

diameter bars each.

(3) Molds 13 through 16 also contained two 38.1 mm

_. (1.5 inch) diameter bars. 1

"A+ R" j

_ (4) On Molds 17 through 21, means Airfoil plus Root and "A/F" means Airfoil only. i

*Blades not included were not single crystals, i
O

:b

P674774 P67477-1

Figure 7. SC Blade Cast with Right-Angle Crystal Selector•

The difference in selector effectiveness is predominantly

associated with exothermic process requirements. Specifically, a

selector design must allow sufficient room between the mold base

and its cross members in order to achieve sufficient exothermic

_', briquette p_cking for proper mold preheat• In this aspect, the

_t

_t_W helix selector design was deficient • A secondary factor potentially associated with the effectiveness of the right-angle selector was the additional 20 mm (0.787 in.) length of the starter

block, which may have increased the probability that only grains

with the [001] primary orientation grew into the selector.

$

16
Blades with m_soriented grains were frequently associated with

_ stray grain nucleation in the right-angle starter To minimize

_ this problem, sharp corners and edges were rounded in the wax pat-

_ tern for the starter. Another attempt to inhibit stray grain

_ nucleation was to increase the temperature gradient by giving cross

members of the starter a slight upward angle (approximately

_ i0 degrees) with respect to the mold base.

&.

h 3.1.2 Wax Assembl_ Configuration

-?

Initially, molds that contained between i0 and 16 blades were

- evaluated. The evolved wax assembly arrangement, which was selected for the remainder of the program, contained 12 blades (or

i test bars) arranged in a 254 _ (ten inch) diameter ring (Fig- ,

ure 8). Two blades were attached to each single-crystal starter,

around a central pouring cup for alloy reception. Each pair of

blades was connected at its tip riser with a 6.4 x 9.5 mm (0.25 x

_ 0.375 center inocfh)thisrunner. bar between A feederthefromtwo the bladespour (Figure cup was 9)connecThus ted, super- to the _-I-_ _

_ heated alloy was top fed through the blade sections, flowing down

_ through the blade roots and crystal selector into the starter cham- _

_ bar. Heat is transferred from the alloy throughout this path,

establishing a gradient in the alloy and superheating the inner

mold walls, promoting single-crystal formation and growth.

A riser was added to feed shrinkage in the thin section of the i

4_

blade root platform. Later, a riser was added to the leading edge }

platform as well, to increase the yield of quality parts• i

3.1.3 Ceramic Shell Mold Manufacture

Initial casting trials were conducted with investment shells

of colloidal silica-bonded zircon and alumina shell systems that

are standard at Jetshapes. These shells were built up to an aver- _!

: 17

__.1
ORIGINAL PPG, E iS

OF POOR QUALITY

.4

%

b

,_, Figure 8. General Arrangement of Single-Crystal Turbine Blades

and Erosion Bars within Exothermically Heated Ceramic

_ Mold.

i[r-_' '] ' "'.%M_, _' 1

POT3M-44

_" Figure 9. Wax Pattern and Ceramic Shell Mold for SC Blades.

_ 18
'7

age thickness of 6 4 to 9.5 mm (0 25 to 0 375 inch) Neither shell

system showed any appreciable tendency toward excessive casting

i inclusions or mold deformation.

Based on the initial casting trials (which compared alumina

shell results with results of a zircon shell) and exothermic pro-

:: cess considerations, the alumina shell was selected for the mold

- !

,.. material. The alumina shell has better fired strength, improved i

._ refractoriness, and better thermal conductivity for transfer of

_. heat from the briquette source to the inner mold walls. The

single-crystal alloy derivatives of Mar-M 247 require the lowest

-_ possible zirconium levels, which points further to the use of

| -:

alumina rather than zirconium silicate to prevent mass transfer to

"_ zirconium from mold to alloy.

_ All molds were autoclave dewaxed, fired at 982°C (1800°F) for

_ _o

one hour, inspected, cleaned, and covered to prevent contamination. '

F

_ 3.2 Exothermic Heating Improvements

i_ In preparation for casting, the fired mold was placed on a _

_ clean refractory ceramic plate in a mold handling frame. A tapered

refractory fiber sleeve was placed around the mold (Figure i0).

I Exothermic nuggets compacted from aluminum and iron oxide powders _

were placed around the mold configuration to ensure heating all !:

parts of the shell. When completely packed, the mold pour cup

covering was removed, and the cup was vacuum cleaned and covered

with a mullite ceramic cover. The mold was then ready for preheating.

Evolution of the exothermic pack process resulted in use of a _!

mold preheat oven for exothermic ignition and firing. Preheat tem- "i

peratures between i010 ° and I093"C (1850 ° and 2000°F) were evalu- i

ated. Best results were obtained using an oven at I093°C (2000°F) i

with an excess oxygen atmosphere. In addition to igniting the exo- _i

19

l !
ORIGil_hL L-i;'.GE b;_

OF POORQUALITY

ji

1 P67368-26

1 MOLD PACKED WITH EXOTHERMIC

i

! BR IQUETS

:'I Fj.c_ure 10. Exo_:herm].c Mo'lcl.

_-0
i

OF POOR QUALITY

thermic reaction between the aluminum and iron oxide, preheating

, reduced mold stresses during the exothermic burn. Ignition and

burn in the preheat oven also eliminated problems with the slight

amount of smoke generated by the exothermic reaction (Figure ii).

Ignition time varied from 5 to 15 minutes, depending on the condition of the nuggets. Burning was completed after another 5 to i0

minute period ]

Mold temperature at the conclusion of the burn was measured to

be approximately 1593°C (2900"F). At this point, the mold was

transferred to the casting furnace (Figure 12).

Variations were noted using the exothermic nuggets; e.g., maximum achieved temperatures varied from 1538 ° to 1732°C (2800" to

3150aF). Single crystals resulted from both extremes, indicating ' _

J

_" , _'_lrflr'_

Figure II. Exothermic Burn in Preheat Furnace.
ORIGINALFhCE I5;

OF poOR QUALI1W,

"" "_ _ :.

L

t It

_t

b ""

i=,lloeo6

REMOVAL OF LOWER SUPPORT

PRIOR TO PLACEMENTIN THE

VACUUM FURNACE

Figure 12. Exothermic Mold Transfer.

22
an adequate heating of the mold even though the indicated temperature reading failed to positively identify the true mold temperature. Further, over a period of time using many different batches

of exothermics, a zinc oxide contaminant (included in the iron ore

used in exothermic preparation) was noted and continued after the

supplier of the briquets changed to a different grade of ore. To

avoid deleterious effect on the casting process (rapid gas evolu-

: tion causing mold runout hal been associated with zinc oxide deposition), the casting production phase used an oxyacetylene torching

• of the mold base after the exothermic burn as s ndard practice.

This torch application eliminates the gas forming coating and

reduces chances for mold runout.

!

• Exothermic materials were purchased to a specification of

: " 1788°C +28°C (3250°F ±50°F) burn temperature With either zircon

_ _ or alumina shells, the molds never indicated an over-temperature

" ,, condition (fusion) during the program. This indicates that for the

_ Jetshapes mold-size sleeve-size preheat system with exothermic

_"_ briquette size and packing, heat transfer to the mold is poor

%:

Much of the heat generated is ineffective in accomplishing the pur-

:, _se, and large variations between exothermic briquette and indi- :

ca_ed temperatures are to be expected•

: 3.3 Castin_ Parameters Optimization _; i

After the mold had been preheated by the exothermic reaction 4

(burn), the mold w[s transferred to the chill plate of the casting :

%

__ unit. Alloy melting was sequenced to the burn transfer mold chamber pump down _o be close to required temperature as the mold was ,_

'i

raised to pouring position. Metal pouring temperatures were deter- "_

i mined by using a two-color pyrometer• Metal was typically poured _'

within 4 to 7 minutes after the mold was removed from the preheat '_

furnace. ._

, 23 1
For the Mar-M 247 alloy, superheats (temperature above the

freezing plateau, P) are considerably higl,er for directional-

: solidification P+236°C (P+425°F) and single-crystal castings

_ P+278Oc (P+500°F) relative to conventional equiaxed castings

i' P+IIIOC (P+200°F). The optimum superheats for casting SC alloy

derivatives of Mar-M 247 were still higher at P+320 ° to 360°C (P+575 °

.. to 650°F). High pour temperature for SC blades were used to inhi- %:

• t bit grain nucleation on mold walls. No significant investment

_ mold-superalloy reactions were observed as a result of the higher

: superheats required for SC castings.

Initial casting trials used Mar-M 247 as the alloy for experimentation. Melting was accomplished in a zirconium-oxide crucible 2

and, as the alloy contains zirconium as an alloying element, no _:

_ alloy contamination was noted. However, higher temperatures were

_ found to be required for derivative SC alloys that had significantly reduced B, Zr, Hf, and C contents relative to the parent

_ Mar-M 247 alloy. The higher temperature and vacuum conditions _ ,

adversely affected the magnesia-stabilized zirconia crucible dur- -

ing melting of the derivative SC alloys. Partial decomposition of

the zirconia crucible resulted in an undesirable addition (up to

0.5 percent) of zirconium to the alloy, which adversely affected

heat treatability and mechanical properties (discussed in Section

5.0).

These results are consistent with the work of Barrett and

3

Lowell that was conducted at NASA-Lewis Research Center.

Zirconium-free NiCrAI alloys were melted in zirconia crucibles and

_.. reported zirconium pickups of up to 0.6-percent weight. (Zirconium

._ additions to their alloys were beneficial from an oxidation-

_ Lcclstance standpoint.)

Melting trials with both MgO and A1203 crucibles gave acceptable results and further blade castings were made using MgO cruci-

: 24

q
J# f

-7

,: bles, which have superior thermal shock resistance. Magnesium

pickup in the SC alloys was 2 to 12 pl_, which was innocuous.

; The technique for the pouring operation proved difficult for

the Jetshapes furnace operator to master. A fast pour could not be

accomplished due to the position of the pour cup in the center of

_-

': the mold. Enlargement 6f the cup did not help. Finally, with a

._" change of operators, casting of the charge within a desired time

. lapse (i second) was accomplished. The more rapid pour lessened "_

alloy heat loss and gave a more uniform filling of the mold. The ,_ j,

more rapid filling of the mold cavity prevented differential tem- _

• peratures within the mold cavity. Results were dramatic: an

• approximate 20-percent increase in the number of single-crystal -_

_, blades generated.

%..

3 4 Post-Castin 9 Treatment

All molds were allowed to cool overnight before shakeout. _

Separation of castings from mold and clinker was carefully done to ,j

. avoid blade surface stresses. After knock-out, the blades were _

salt bath cleaned and acid etched for preliminary grain evaluation. ;

i Blades accepted by the first inspection were solution heat

treated to specification. They were then returned to Jetshapes for ;_

2_

sand blast and grain etch and X-ray diffraction (XRD) evaluation of •

"_ the crystal orientation. Acceptable blades then wece processed

4

through the usual foundry finishing operation: visual, fluorescent ._

penetrant, and x-ray inspections, i_

i

25
SECTION IV

4.0 SC MAR-M 247 PROPEPTY EVALUATION

In addition to adapt£ng the exothermic process to the manufacture of SC turbine components, the program characterized primary

mechanical properties of SC Mar-M 247, providing a baseline for

subsequent development of derivative alloys with improved tempera-

: ture capabilities. Tensile tests were performed at room temper-

' ature and 760°C (1400°F). Stress-rupture data were obtained for SC

Mar-M 247 at 760°C (1400°F), 982"C (1800°F), I038°C (1900°F), and

1093°C (2000"F). In all cases, the test plan involved both longitudinal and transverse tests from heat-treated blades. Test

specimen machining, tensile testing, and rupture testing were per-

": formed at Joliet Metallurgical Laboratories (Joliet, Illinois).

0 Blades were evaluated with two different solution heat treat-

_; ment temperatures: 1232°C (2250°F) and 1246°C (2275°F). The first

_ solution heat treatment was the recommended solutioning treatment

1 +'-':_

for DS Mar-M 247 from Project i. The second (slightly higher) : _

, solution treatment was also evaluated in Project 1 and resulted in

: better homogenization of the matrix, slightly improved average rup-

+ ture lives, and slightly increased amounts of incipient melting.

: ..'

I

Most of the testing was performed with 1.8 mm (0.070 inch) .

t gauge diameter minibars machined from SC blades with known orienta-

![ tions. Some testing was also performed with 0 " 51 mm (0 . 020 inch)

! thick flat specimens machined from the airfoil in order to obtain a

quantitative assessment of thln-wall effects on stress-rupture properties. Machined-from-blade (MFB) specimens an4 their relationship to the MATE SC blade are shown in Figure 13.

In addition to the MFB specimens, several 1.8 mm (0.070 in.)

t

¢ diameter mlnibars were also machined with a transverse orientation {

! from small 38 mm (1.496 in.) diameter SC ingots. Ii+

i PRECEDING PAGE BLANK NOT FILMEI_

1 _mTENrlOm t BL K.Ly + !I

|

i
ORIGINAL PAGE IS

OF POOR QUALITY

oh---,

r---1

i i

i i --

I

• " _.1

MATE SC TURBINEBLAOE PTS641

<-..

540uqcTHREAO

P3.19(IL125) DIA. r-l.solo.oTo) DIA. \

', ,1,1IlmIil • - (0,221-iD ,m.,--,---- 10.1610,401-_

: I_ 25,40(tOOl

NOTE: ALL DIMENSIONSIN MM {IN.)

Figure 13. MFB Tensile and Stress-Rupture Bars.

28

.•w_ ®
4.1 Tensile Tests

Room temperature and 760°C (1400 °F) tensile test data for

specimens machined from SC Mar-M 247 turbine blades are provided in

Table 2 and 3. For single crystals with the [001] orientation,

average tensile strengths were slightly lower and ductility was

: slightly higher relative to multigrained DS Mar-M 247. In contrast

to stress-rupture results, 0.5 mm (0.020 inch) thick flat specimens

:_ exhibited slightly higher strengths than 1.8 nun (0.070 inch) diameter round specimens. Transverse single-crystals exhibited average

tensile strengths and ductilities that were slightly higher than

those of DS Mar-M 247.

.-

L

Yield strength data for longitudinal MFB specimens and similar

. data for 6.35 mm (0.248 in.) diameter SC Mar-M 247 specimens gener-

•_j ated by MacKay, Dreshfield, and Maier 4 are presented graphically in

'" Figure 10. Substantial variation, indicated by a scatter band

,,: width of -+82.7 MPa (-+12 ksi), is present in these data. No obvious _,.

J - ._

strength trends were present in the temperature--over the range of

i 23 ° to 760°C (73° to 1400"F)--or specimen section size. In fact, :

c

the strengths of the 0.5 mm (0.020 in.) thick flat specimens were

• unexpectedly higher on the average relative to the 1.8 (0.070) and

6.35 mm (0.248 in.) diameter specimens.

4

Inspection of Figure 14 also indicates that the yield strength

: is reduced by approximately 6.9 MPa (1 ksi) for each degree of

I deviation between the [001] crystallographic direction and the _

blade axis. MacKay, et al., 4 indicate that this decrease is asso- ',

ciated with the increase in the Schmid factor for the 1111) [I011

slip system as the specimen deviates from the [0011.

Several tests were conducted at 760°C (1400°F) for orientations near the [210] and [311] specimen axes. The data in Table 3

_-_ indicates that these non-[O01] orientations had significantly lower

_" tensile strengths.

29
ORIGINAL PAGE IS

OF POOR QUALITY

5

Z _

_ ,,oeoe,.l

<

_ _ o

eeleaeeeeeeeele • • • *eemeii•mee•*•* _ _

• _ .......... • BBBB °oo .....oo

illli| |l||lll| _ illlili_|||||_| _

000 000_00000 0

0

_0000 0

_ 0 O0 O_ _ _ _ O_

oO _ '_>>_>_ "_ _ _ o

0.0*

' "-" II |®_

Jo_ _ _

N lllllllllllllll M M . ,lll?_ll,lll,_ _ _ _ II

30
k_'_ ' C

ORIGINALPAGE|g

OF POOR QUALITY.

_ _ -

• o,,.o*.ee *J,.oi.ee..o.oe

"'" _ 0

• _ - ................. ._ _ _ _ _g_J_gd_dd_ _

: o !

- _ lll|lllll||llli I -"

0

31
ORIGINAL PAGE IS

OF POOR QUALITY

1103(IN) .. 100

10,34 _" 150

-- _'_-_"_'_-,,,,_ I 14o.o

0.(,3o, - _ ,30

.

027 11201' _ _ 120

SYMBOL SPECIMEN TEMPERATURE

_3 1.0mm(0.070 INCH) OlA MFB 23°G (73°F)

O 0.5turn(0.02 INCH) FLAT MFB 23°C (73°F) ._

- I.Bmm(0.O70 INCH) DIA MFB 7BO°C (1400°F) 110

" - 758 (110) -- O [3 0.5ram(0.02 INCH) FLAT MFD 780°C (14DO°F)

:'" L_ 0.3mm(0.26 INCH) OIA 23°C (73°F)

<_ _7 6.3ram(0.25 INCH) OIA 6411°C(1200°F) ;" -- 0.3mm(0.26 INCH) DIA 760°C (1400°F)

690 (100)

1DO

6 I0 15

ANGLE BETWEENBLADE AXIS AND IDOl], OEGREEB , ,._.

Figure 14. Yield Strength of SC Mar-M 247. "

The transverse properties, where orientations were not controlled, are almost identical to average DS blade properties. The

two low -alues at 760°C (1400°F) strongly suggest weak secondary

orientations, when compared to the known orientations of the longitudinal spec linens.

..

Elastic moduli were calculated from load-elongation curves for

longitudinal and transverse orientation tensile specimens tested at

' room temperature and 760°C (1400°F). However, since the extenso-

' metry was attached to the grips [that is, not on the 1.8 mm (0.070

in.) diameter test specimen], the calculated values were not considered reliable and are not reported. (For a given crystallographic orientation, calculated modulus values ranged from equal to

well below those reported in the literature for pure nickel and

Mar-M 200.)

32

I

1
Transverse tensile data for controlled orientation 1.8 mm

(0.070 in.) diameter minibars machined from SC Mar-M 247 ingots are

provided in Table 4. This data exhibited severe scatter for the

[100] orientation, but little scatter for data taken 15 ° from the

[100]. Ignoring the two low data points at [100] gives yield

strengths comparable to the longitudinal data in Figure 14.

." 4.2 Stress-Rupture Tests

Stress-rupture testing of SC Mar-M 247 was conducted in the

range of 760 ° to 1093°C (1400 ° to 2000°F). At 760°C, data was

obtained using 1.8 mm (0.070 inch) diameter specimens machined from

blades (Table 5) and ingots (Table 6). Both MFB 1.8 mm (0.070

inch) diameter minibars and 0.5 mm (0.020 inch) thick flat speci-

>.

mens were used to obtain stress-rupture data at 982 ° (1800°), 1038 °

.27

: (1900 ° ) and I093°C (2000°F); these data are provided in Tables 7,

¢ 8, and 9, respectively.

2 +

As expected, the greatest variation in results was observed at I'+_

760°C (1400°F), where rupture life has been shown to be highly

dependent on SC orientation.

£

For specimen orientations within i0 ° of the [001], substantial

scatter was observed. This variability in rupture life is similar

to that reported for SC Mar-M 247 by MacKay et al.; 4 they concluded

1

_. that the amount of crystallographic rotation associate_ with +

.i

primary c_eep on the IIIi} <112> slip system could explai- the variability _n rupture life. +

._ The effect of solutioning temperature on 760°C (1400_'F)

i stress-rupture lives was also noticeable in the [001] orientatzon

longitu_!nal specimen data; i.e., specimens solutioned at 1246°C

had average rupture lives that were 40 percent greater 

than those of specimens solutioned at 1232"C (2250°F).

33
TABLE 4. TENSILE TEST RESULTS ON 1.8 MM (0.070 INCH) DIAMETER

CONTROLLED TRANSVERSE ORIENTATION MINIBAR5 MACHINED FROM

SINGLE-CRYSTAL INGOTS OF HEAT-TREATED* MAR-M 247.

Ultimate 0.2%

Tensile Yield

Specimen Orientation ITemperature Strencth Strength Elongation R of A

Number of Specimen °C (°F) MPa (ksl) MPa (ksi) (%) (%)

397-21S [i00] RT 834 (121) 821 (i19) 10.6 27.9

398-21S [100] RT 1124 (163) 917 (133) 10.4 20.0

397-23S 15" to [100]** RT 834 (121) 814 (i18) 6.1 17.6

398-31S 15 ° to [100]** RT 876 (127) 841 (122) 10.6 18.0

397-11S [110] RT 896 (130) 889 (129) 9.0 17.6

398-11S [110] RT 883 (128) 869 (126) 12.1 30.8

397-22S [100] 760 (1400) 834 (121) 793 (115) 8.4 17.6

398-22S [I00] 760 (1400) 1193 (173) _93 (144) 8.2 13.9

398-32S 15" to [i00]** 760 (1400) 931 (135) 827 (120) 11.8 18.0 _,

398-34S 15 ° to [i00]** 760 (1400) 88? (129) 814 (118) 9.0 17.0

397-12S [II0] 760 (1400) 972 (141) 827 (120) 11.7 13.0

398-12S [110] 760 (1400) 848 (123) 807 (117) 14.3 27.6

*Heat-treated at 12460C (2275°F) for 2 hours, plus 982°C (1800°F) for 5 hcurs,

plus 871°C (1600°F) fo_ 20 hours. _"

*'15" rotation away from [i00] toward the [ii0].

Limited data for SC specimens with orientations near the [311]

and [Iii] indicated that these orientations were stronger than the

[001]. The high strength of the [iii] specimen was expected based

on previous studies with SC Mar-M 200 and Mar-M 247.4'5 On the

other hand, MacKay had reported the [311] and [331] to be relatively weak.

As anticipated, based on SC Mar-M 200 and Mar-M 247 studies,

single crystals with orientations near the [ii0] had very low

stress-rupture life. This result was attributed to extensive pri4

mary creep on a single Illl} <112> slip system.

Transverse minibar stress-rupture testing at 760"C (1400°F) on

specimens machined from the SC ingots indicated that specimens with

an axis rotated 15 degrees away from the [001J (rotated toward the

34

i
OF POOR QUALITY

TABLE 5. 760Oc (1400°F) STRESS-RUPTURE TEST RESULTS ON 1.8 Ml4 (0.070 INCH) DIAMETER

BAMS MACHINED FROM SINGLE-CRYSTAL TFE731-MATE TURBINE BLADES OF HEAT-TREATED*

MAR-M 247.

i_oi __o I i L

CEyStal Treatment Rupture

ISpeclmen Selector orientation TempeEature Stresm Ll[e Elon9atlon H of A

I Number TYPee* Of Specimen *C (*F_ MPe [ksz) [Hours} {6_ [|_ . kelirki

LdNGZTUDIHAL $PECII,_NS

r s04-ZL 4, 0.5" Off [1003 1232 (2250) 724 (1053 - J 3.0 7.3 FaiZe_ on loading th(ough

8hr;nk defect.

303-1L 3n 6" off [001] 1232 (2250) 724 (105) 78.2 4.4 16.6

:294-1L 1H 10" off [001) 1232 (2250) 724 (105) 165.3 15,6 19.3

295-1L 2B 5* Off [001J 1232 (2250) 724 (1053 54.5 0.5 25.3

!295-8L 2B 4" Off [001) 1232 (22503 724 (1053 13.8 14.2 19.8

295-2L 2H 4" _ff [311] 1232 (2250) 724 (105_ 496.5 4.7 14.2

1295-10L 2H 28" off [0011 1232 (22503 724 (105) 31.7 11.5 21.6

295-3L 2H 7* Off [110| 1232 (2250) 724 (105) 26.9 5.9 10.2

295-13L 2H 6* Off [3113 1232 (2250) 724 (105) 214.8 9.2 15.3

1294-6L IH 6" off [111[ 1232 (2250) 724 (105) 225.1 9.5 IJ.9

405-3L 18R Vllual [001| 1246 (2275) 724 (105) 245.6 13.4 16.3

1405-4L 18R Visual [001[ 1246 (2275) 724 (105) 36.0 5.9 12.1

406-2L 19R Vleuel [001] 1246 (2275) _24 (105) 66.9 11.9 25.5 Failed through

]406-6L 19R V|eual 1001] 1246 (2275) 724 (105) 82.3 10.3 17.6 shrink

408-9L 21R Vieual 100!] 1246_75) 724 (ld53 119.2 9.1 16.1

., (AnticiLxlted Life - DS Caltin_s 60, to 300 Hours i

TRANSVE I_E SPECIMENS

rS03-1T 3R *** 1232 (22503 655 [95) 276.7' 8.5 "' 21.2 "'

!294-1T I_ 10" nil [Ii0] 1232 (2250) 655 (95) 1.0 9.3 12.6

•. 295-1T 2H *** 1232 (2250) 655 (95) 348.8 4.4 15.4

295-8T 2H e.* 1232 (2250) 655 (953 20.4 2.8 10.4 Failed thtouqh

ehrtnk

:_. 295-2T 2H 4" off 1331] 1232 (2250) 6 5 (95) 759.5 11.3 27.0 295-IOT 2H **" 1232 (2250) 6.5 (95) 60.1 12.7 17.0 Patled through

• shrink

_*, 295-3T 2H 2* Off [110] 1232 (22503 655 (95) 2.8 16.4 42.4

_" 295-13T 2H *** 1232 (2250) 655 (95) 501.8 3.4 10.4

.," 294-6T IH 9" off 1210[ 1232 (2250) 655 (95) 48.8 6.5 20.8

?, 405-1T 18R Visual [100] 1266 (2275) 655 (95) 1124.5 11.6 21.1

405-4T 18R Vlaual [100) 4246 (2275) 655 (95) 598.2 10.7 20.5

406-2T 19R VLsual [100J 1246 (2275) 655 (95) 928.4 11.7 22.7

406-6T 19R V/sum1 [1001 1246 (2275) 655 (95) 849.3 16.3 20.2

408-9T 21R Vlsual [I00) 1246 (2275) 655 (95) 571.2 15.1 27.0

• 40P-I1T 21R Visual I100J 1246 {2275} I 655 _953 1145.3 11 1 18.0

{Ant|cipate_ L|fe - DS Castlngl 10_rs I ; ,_

q_-_olution t_'eated as |r_lcated for 2 hou[e with Argon Quench, plus 982 C (1800 F) fo_ 5 hours, |

" plus 871"C (1600"F) for 20 hours. t

• * R • Right-angle crystal 8eleCtorl H = Helix crystal selector. _,

), ***Transverse orientatione not determined.

%

[ii0]) had comparable lit.es to [001] specimens (Table 6). This

result is in agreement with MacKay's results. 4

In agreement with the MFB data and studies cited earlier,

machined-from-ingot specimens with the [110] orientation had rupture lives that were two orders of magnitude lower than those near

* the [001] . (

• e

i '

4-'? '

As anticipated from previous studies, there was far less

: scatter in the higher temperature range 982 ° to I093°C (1800 ° to

2000OF) stress-rupture lives. (See Tabl_s 7, 8, and 9.) At these

temperatures, the results for [001] oriented specimens were found

J

4

35 _
ORIGINAL P;" _' !_-:

OF PoOR Q,.._,",,'..,_i,

TABLE 6. 7600C (1400°F) _TRESS-RUPTURE TEST RESULTS ON 1.8 MM

(0.070 INCH) DIAmeTER CONTROLLED TRANSVERSE ORIENTATION MINIBARS MA_INED FROM SINGLE-CRYSTAL INGOTS OF

HEAT-TREATED e MAR-M 247.

Rupture I

Specimen Orientation Strtsa Llfe Elongation IRofA

Number of Specimen NIPs (ksi) (Hours) (t) I(t)

I

3ge-23s [100] 724 (103) 227.1 10.2 ) 20.2

i

398-24S [100] 724 (105) 186,0 13.2 17.6

i

398-33S 15e off [I00] 724 (i0._,) 145.3 13.0 16.7

398-35S 15 e off [100] 724 (105) 356.3 9.2 18.7

398-13S |110] 724 (105) 1.3 28.8 37.8

398-14S [ii0] 724 (105) 1.6 30.7 36.1 r

i

*Heat-treated at 1246eC (2275eF) for 2 ho_rs with Argon Quench, [

plus 902°C (1800°F) for 5 hours, plus 871"C (1600°FI for I

20 hours, i

to be at the high end of previously documented DS Har-H 247 blade

results from MATE Project I. Scatter in the transverse specimens

was dramatically reduced.

Similar results have been observed with the SC Mar-M 200

alloy. 6'7 In those studies, it was noted that for 857°C (1575°F)

and higher temperatures, primary creep (which is associated with

deformation by a single slip system) is reduced by an order of magnitude below that observed at 760°C (1400°F).

Examination of 982°C (1800°F) longitudinal and transverse SC

Mar-M 247 stress-rupture data (Table 7) on <100> oriented crystals

indicates that solutioning at 1246°C (2275°F) yields a small

improvement in average stress-rupture life; i.e., 6 percent for

: longitudinal (i00) and 22 percent for transverse material.

• i

Inspection of I038°C (1900°F) stress-rupture data (Table 8)

also indicates that the higher solutioning temperature resulted in

improved rupture lives. The average improvement in longitudinal

properties was marginal (4 percent), but the transverse data exhibited a 44-percent improvement.

36

®
ORIGINAL P,_QE t_

OF POOR QUALITY

TABLE 7. STRESS-RUPTURE TEST RESULTS AT 982"C (1800°F) ON TEST SPECIMENS MACHINED

FROM SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED* HAR-M 247.

Mold Iio. _lut Ion

Crystel Triatmnt Rupture

g_clmin gila|or Orlentltlon k_returt I_clMn ItrelI Life II_qatlon R oI A

N_r TYPe it O{ H_CJBn *C(*T) TypQ MFI (kll) (Houri) (4) (I) llotlrkl

_ngltud|nll H_¢lBnl

304-21 48 i" off [o01] 1232 (2250) 1.8 m 207 (Jo) $4.1_ .... Nllid tr_

B m

(?0 el}) r_nd

303-21 3R S* Off (00l| 10O.S 18.t S0.1 Adapter

295-111 2H 8" oft (001] 110.0 17.1 50.1

295-141 28 4" Oft [001] 90.0 28.0 53.6

]03-SL 3H E.5" off {001] 103.7 25.0 49.7

]03-8L ]H 4" off (001| I 114.7 21.2 47.4 295-4b 2R 8" off [OOl] I u 123.8 21.1 48.9

295-71 2H 8" Off [211| j 102.8 14.2 32.7

• 29_-91 2e 6" off (311J l i 240.4 15,1 31.0

322-13L SH 4" off (O01] 1246 (2275) ; i 102.3 21.0 48.5 t

12}-121 4R 1" O_f (001| 113.4 20.4 47.4

32]-141 6R 3" off (001| _ ; 124.5 17.3 44.4

322-13F 5R 6" off [001) 0.5 mm 02.4 36.9 --

123-12r IH 1" off [OOl] (20 all) flit I 149.3 34.0 --

323-14P 6R 3" ore [OOl1 J 1 49.2 ,3.1 --

/

(MtJell_tid lif0 - 1.8 mm (70 111) r_ndl frl 09 Castles1 50-120 _urs)

t (_tJoJ mild 11fl - 0.5 Im (20 Ill) flare fr_ DB Caitinqml 1S-)0 _ura)

Trinivorle 8_cJlHInl

[104-2T r ' 4H' _t 1232 (2250) 1.0 u 1 106 (27) 1 120.0 ] 17.1 ] 28.2

| I _terminid • (70 m,l_ roun_m I m [ I

1.}-. ; 3H , r 0 m i / z..2 I 19.1I 38._

2. I / I / I t21"'5 22 , ,4 .hr,o |19_-4:r 2. i l r , i / 183.3i 17._ I 48.s

_ Ihrlnk

- T 1246 22751 . . .

" /34_-4T I 7. i ! / / 129_1 ! 10_9 _ 44:0

/}(343-644- T I}v,. I I I ) i / 229.7 J 17.) ( 4}.8

. I I I l m241.]1 200 i47.s

' i3,8., 1,H i , , i2o,,30m3_o,2 1 191 i}8,

398-9T 14R | | | I | 111.7 | 24.5 _ 44.5

" (404-sT ,7, l I ' I m i0,.8i ,.s 27.8 ,,04-. _7, I s I m | m_0,1 | ., _,5

-" F40_-4z J 2o, / 1 / I / i 87.6 I 27.7 I 43.6

* _ L_ t v V0lii l I

6' MP 27 k|i

_ Anticipated lift fr_ DS Caetings_ I 0 Routs t 18 S ( )

$.__]*S°lutIontreated as Indlc.ted with Argon Quench, plus 982"C (ZS00"r) for 5 hours, pius 871"C (IG00*F) for 20 hours. ),

[ **R " Rtght-A_le crystal leIectorf H - Helix cryitsl lelector. '_

i

Average 982 ° and i038°C (1800 ° and 1900°F) rupture lives of

0.5 _ (0.020 inch) thick flat specimens [solutioned at 1246°C

; (2275°F)] are, respectively, 82 and 89 percent of the average lives

of 1.8 _ (0.070 inch) diameter specimens solutioned at the same

temperature; i.e., rupture lives of the thinner specimens are surprisingly high at these temperatures relative to equiaxed casting

section-size trends. Retention of stress-rupture life in the

0.5 mm (0.020 in.) flat specimens apparently reflects the absence

of grain boundaries as weak links that cause early failures. In

contrast, at i093°C (2000°F), the 0.5 mm (0.020 in.) thick flat

specimens have only 20 percent of the rupture life of the 1.8 mm

, (70 mil) diameter specimens (see Table 9). This difference is partially attributable to the increasing effects of oxidation (reduc- ._

ing the effective section size) as well as the section-size effect. :'

37
TABLE 8. STRESS-RUP?URE TEST £ESULTS AT 1038°C (1900°_) ON TEST SPECI_'ENS MACHINED

FROM SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED* MAR-M 247.

Nold NO. _ qoZut/on

CrygteZ Tr.._nt Rupture

• l_¢Imen I SeleCtor Orisntation _mptrature 8peo_en Stren8 Life Elm_letion • of A

Hustle Type H of •peoLnen *C('71 Type NPl (klI) (Hourn) (41 (t) Rlm_kl

. .

ten, It ud ins1 8pmoimens

. 2_5-12_ 3n 9" o_f fOOl( 12_2 (2350) 1 1.s m 131 (191 228.3 1%0 49.7

39_-61 2B l* off foo,f (?o ILl) round 193.0 21.6 S0,3

303-3_ 3R 4_ off lOlL| 312.0 23,2 58.2

303-9L 3R 3 ° Oft f001] 141.3 30.6 62.3

303-7L 3R 9.5 _ off [001] 179.6 24.7 48.3

303-111 3• 2.5" off [001f _ / 234.0 23.6 51.9

323-101 6R 2* elf f001] 1246 (22/51 1 207.? 17.8 54.7 ]44-9L 4• 2 ° Oft 1001] 205.9 32.1 58.9

0.5 nee 184.3 13.4 --

323-10r 6• 2 ° Off f001| (20 nIll flat

(AnticIpoted II_e-DS Cnotlng|! 55-160 hourm)

Transver i gpecimens

.... - _ . . _ . .

295-6T 28 Mot Determined 1232 (2250) 1.8 mm 131 (191 117.7 14.) 37.S

3C3-3T _ • (70 mill round 121.4 10.9 35.5

!30_-9T 3• _ 47.9 6.0 16.7 railed through

! shrink

323-13T S• 1246 (2275) 231.8 21._ 50.9

323-15T 6• 209.8 20.9 50.9

344-4_ 8• 132.1 15.4 33.1

344-_T 8• 112.7 17.9 53.6 ,_

398-ST 14• 230.8 27.9 48.6

_- 407-4T 20• F 115.9 . ! 15.9 43.9 I

(kntlolp_ted life - U• Csstln9|_ 100 Ho_rml

• Solution treated us Indicated _lth Argon 0ueneh. plul 902*C (1000*_1 _or 5 hours, plus 871*C (1600°r) for 20 hours.

_ **A = Right-Angle crystal selector! H - Hells rryn_sl selector.

TABLE 9. STRESS-RUPTURE TEST RESULTS AT I093°C (2000°F) ON TES- SPECIMENS MACHINED

" FROM SINGLE-CRYSTAL TURBINE BLADES OF HE_T-TREATED' _-M 247.

Mold No. •olutLon [

Ceyatsl _restJNnt I •upturn

Specimen Selectur Orientation ?amp°rotors I Specimen •trees LI_a •longstio., R o_ A

NUlit_lr Type** Of •pecimen *C(*P) Type NPs (ksI) (Hours) (t) (t) •eoarkn

IX_ng i t od innl Specimens

327-151 4• 9* of_ f001] (70 I111 round 103 llSI 60.0 21.5 59.1

343-4_ 7• 3 ° off 1001] 103 (15) 39.4 13.7 33.6

406-11 19• Visual f001| 97 (141 ...... qechtne Nsl-

• _unetion

) 323-13r 4• 10" oft f001f 0.5 am 103 f15) _.5 24.0 --

323-15r_ 4• 9 • ot_ f001| (20 nil) fist 103 (15) 9.3 2b.l --

(Antl©.peted life - 1.8 m (70 mll} rc _ndl _,'m D• CeotLngsJ 35 hours at 103 14Pa (15 kml)

_398-3T 14R Not Determined 1246 (23151 1.| mm )? (14) 03.3 32.7 40.4

!390-12T 14• (70 mill round J 101.3 12.0 17.0

17• J q 43._ 16.1 46.7

404-3T t - I _ 1

1 404-1T 19• / 66.9 24.8 39.4 407*IT 20• 76,4 20.7 38.5

_ 400-4T 21R 71,4 16.4 49.4

(AnticiPated life - 1.0 m (70 re;l) ro_nd• from Dg Ceeti_gu_ 53 bourn at _1 N•m (li tail

• Solution treated as indicated vith Argon Quench, plus SO2*C (lJ00*_) tor Shears, plue 0110c (14000p) _or 20he.Jr°.

• *• . Right-angle orystel _ele_tor! • • Hells ©rystaL selector.

38
SECTION V

5.0 SC ALLOY DEVELOPMENT

Previous experience with SC Mar-M 200 and the work conducted

in this program with SC Mar-M 247 demonstrated that the [001]

orientation SC alloys were not significantly stronger than the DS

form in stress rupture. Consequently, it was necess_y to exploit

the alloying freedom provided by the SC alloy macrostructure to

obtain a significant improvement in mechanical properties.

For the Mar-M 247 alloy, the obvious modification was to

remove all (or part) of the elements (such as B, Zr, Hf, and C) that

strengthen grain boundaries in conventionally cast and DS alloys.

Since B, Zr, and Hf are also strong melting point depressants,

their removal permits the eutectic gamma prime, which does not contribute significantly to strength, to be more fully solutioned and

" then reprecipita_ed as fine gamma prime. This approach has also

9 and Gell I0 ii for other been successfully applied by Boone, 8 Duhl, ' .

SC nickel-base superalloys.

12

Based on the work of Jackson, stress-rupture life is a

strong function of the volume fraction of fine _' in the microstructure. Gamma prime in Mar-M 247--even when heat-treated at (or

near) the incipient melting temperature--is only partially solutioned; i.e., substantial amounts of eutectic T' remain. Thus, the

melting point limitation of the unmodified Mar-M 247 inhibits

realization of the full strengthening potential of the 7' in this

alloy.

_ In addition, cobalt removal or reduction to levels lower than r

: the i0 percent Co present in Mar-M 247 was considered to reduce

alloy cost.

Some of the grain boundary strengthening elements such as Hf

also enhance oxide scale adhesion on diffusion aluminide coatings.

39
Consequently, in some instances, a small amount of Hf (or Y, which

also enhances oxide scale adhesion) was retained for potential oxidation resistance benefits.

The approach to alloy modification was performed in three

steps:

o An analytical study of eight "paper alloys"

o A microstructural stability study with equiaxed derivatives

o Development and characterization of SC derivatives of

Mar-M 247.

These aspects of SC alloy development are reviewed in the following

: paragraphs.

5.1 "Paper Alloy" Study ./__

Based on the above considerations, a preliminary analytical

study was performed on eight "paper alloys" derived from the

Mar-M 247 chemistry. Alloy chemistries are provided in Table i0. -

The analysis consisted of computer calculations of electron

vacancy numbers using both NV3B and NVss methods. Since both

yielded the same results once carbon and boron were removed, only

results of the NVss calculations are reported in Table ii. Also

included in Table ii are calculated volume fractions for the phases

|,

? and alloy density.

I

I This study indicated that all of the candidate SC alloy derivatives had lower calculated electron vacancy numbers than the

_ maximum specification composition of Mar-M 247. Therefore, modifi- &

i cations 3 through 8 were assumed to also be metallurgically stable;

I that is, free of any topologically close packed (TCP) phases.

40

l
J

ORIGINAL PP,_:',_;.'_'_

OF POOR Q_.-.,-.:fY

TABLE i0. COMPOSITIONS (PERCENT WEIGHT) OF MAR-M 247 AND DERIVATIVE SC ALLOYS.

compositions of Alloys

..... i

1 2 3 4 5 6 7 8

MW_O ....

"- 0.50 0.80 0.80 0.70 0."0 0.76 0.72 0.76 ;

9.50 10.50 10.50 lO. 20 lO. 50 9.98 9.45 9.93

"_ Ta 2.80 3.30 3.30 3.I0 3.30 3.14 2.97 3.12

Ti 0.90 i.20 i.20 i.I0 I.20 i.14 1.08 i.14 "-

" Cr 8.00 8.80 8.80 8.60 8.80 8.36 7.92 3.32 ,_

.' CO 9.00 10.00 - - 5. O0 10. O0 5. O0 "/

Fe 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15

"..

A1 5.30 5.70 5.70 5.$5 5.70 5.42 5.13 5.39 _.

8 0.010 0.020 .... :_

"" Zr 0.030 0.080 - - -

_.'-, c o.13 o.17 - - - i

_',' Hf 1.20 1.60 - 0.60 - 0.60 ._,

: Hn 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05

' 51 , 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05

! h

._ ! _Ji [ Bal Bal Bal Bal Bal Ba! Bal Bal _4

?: : ........

NOTE: No. 1 is minimum of normal specification range. ;_

No. 2 is maximum of normal speciflcation range. -_

l

• No. 3 zS maximum of range with zero B, C. Hf, Zr, and Co. 4

¢ NO. 4 is mid-range chemistry with zero B, C. Hf, Zr, and Co. _.

No. 5 is alloy NO. 3 with 0.6-percent Hr.

NO. 6 is alloy No. 3 w_th 5.0-percent Co. ;i_

J (

NO. 7 IS alloy No. ] with 10.0-percent Co. 4

. NO. 8 iS alloy No. 3 with 5.0-percent Co and 0.6-percent Hr. :

: ¶_

, m

.f-

"e

!

,: 41
ORIGlr_AL PAGE ;-_

OF POOR QUALITY

TABLE II. CONI_TER ANALYSIS OF NAR-M 247 A_D DERIVATIVE CONPOSITIONS LISTED

IN _ABLE I0.

Modification No. Carbides (V/O) Bo(Vr/

iOde)s (Vy' /O) (V/7O) __MV_s g/ca 3 (1b/In.P3)

1 (,=in.) 2.90 0.14 54.74 42.45 1.93 8.61 (0.311)

2 (lax.) 3.81 0.28 60.57 35.64 2.32 8.64 (0.312)

3 0 0 61.91 38.10 2.13 8.61 (0.311)

4 0 0 59.70 40.30 2.01 8.61 (0. 311)

5 0 0 62.16 37.84 2.14 8.64 (C. 312)

6 0 0 58.83 41.17 2.08 8.64 (0.312)

7 0 0 55.65 44.35 2.03 8.64 (0.312)

8 0 0 J 58.77 ,_1.24 2.07 8.f7 (0.313)

|

5.2 Microstructural Stability Study

To assess the thermal stability of alloy modifications of

Mar-M 247, five chemistries were selected, melted, and cast into -,r

equiaxed ingots. The first composition, Alloy i, is basically

Mar-M 247 with hafnium, zirconium, cobalt, and boron eliminated and

carbon held at an absolute minimum level. This composition was

vacuum melted by Cannon-Muskegon (Muskegon, Michigan) to a very

high purity (Table 12). Other alloy compositional variations were

produced by remelting portions of the resulting master alloy and

adding minor amounts of selected elements during remelting and

casting at Jetshapes.

The aim and actual compositions of the five alloy modifications are listed in Table 13. Alloy 2 restores half the normal

hafnium in Mar-M 247 to the Alloy 1 composition. Alloy 3 returns

half the normal cobalt and hafnium in Mar-M 247 to Alloy i.

Alloy 4 is a molybdenum addition heat, while Alloy 5 is an yttrium

modification for environmental resistance.

42

®
ORIGINAL _

OF POORQUALITY

TABLE 12. CHEMICAL COMPOSITION OF MASTER HEAT NO. VF-31

OF =STRIPPED" MAR-M 247 MELTED BY CANNONMUSKEGON CORPORATION.

' Alloy Composition (Weight Percent)

Element Aim Chemistry Heat VF-3_

•, Ho 0.80 0.8

.... W 10.50 10.2 :

• Ta 3.30 3.4

Ti i. 20 i. 1

Cr 8.80 9.0

t

CO - -

A1 5.70 5.5 3

C - ,_. 018

Hf -

; Zr -

" Oxygen* 4

:>" Nit_ogen* 6

•"_'_ Ag * < 1 :.

".- Bi* < 0.5 _.

,-, Mg* 7

" Pb*. < 1 "

,:_ Se * < I . _"

Te* < 0.5

: " TI* < I _'

Ni Bal y

• Parts per million for gases and trace elements . _.

The solidus and liquidus temperatures of Mar-M 247 and the _

_ five alloy modifications were determined by differential thermal %.._

_. analysis (DTA) through interpretation of the cooling curves. /<

Although the values may not have been exact, they were directly ,_

_ comparable and are presented in Table 14. "_.

'_ To assess alloy sensitivity to sigma and mu phase precipi- ,_

_' tation, stability tests were conducted on the five alloys using !._

exposures of 843°C (1550°F)/i176 hours and I0930C (20000F)/50 "_

_ hours, respectively. Prior to testing, the alloy samples were !_

given either the solutioy_plus coating cycle plus age or the coatt.

i ing cycle plus age heat treatments Two solution treatments were
ORIG!NALP_._. ;_

OF POOR QUALIW

TABLE 13. AIM AND ACTUAL COMPOSITIONS OF MODIF:ED MAR-M 247 ALLOYS.

Alloy Composition (Weight Percent)

Alloy 1 Alloy 2 Alloy 3 Alloy 4 Alloy 5

Ele-,e;tt ALl Actual Aim Actual ALB Actual Aia Actual Aim Actual

No 0.80 0.0 0.8 0.76 0.8 0.77 2.0 1.70 0.8 0.75

w 10.50 10.2 10.2 10.27 9.4 10.03 10.1 10.21 10.2 10.04

Ta 3.30 3.4 3.4 3.25 3.2 3.24 3.4 3.00 3.4 3.07

Ti 1.20 1.1 1.1 1.15 1.0 1.10 1.1 1.10 1.1 1.11

¢r 0.00 9.0 9.0 9.10 8.5 8.67 8.9 8.77 9.0 8.78

Co 0 ...... 5.0 5.13 ........

A1 5.70 5.75 5.75 5.48 5.45 5.42 5.65 5.47 5.75 5.52

C 0 0.010 0.018 0.017 0.018 0.014 0.010 0.015 0.018 0.015

Rf 0 -- 0.6 0.63 0.6 0.60 ........ .%

Zr 0 0 LAP** <0.01 LAP <0.01 ........

- Oxygen* -- 4 ................

Nitrogen* -- 6 ................

Ni Bal Bal Bal Bal Bal Bal Bal Sal Bal Bal

I

i ¥ 0 .............. 0.4 0.50 :_..

*PPM ": " _ " ""

**LOW as possible.

NOTE: Actual composition of Alloy I furnished by Cannon-Muskegon Corp. ;

alt other analyses from National Spectrographic Corp. (Cleveland, Ohio).

TABLE 14, SOLIDUS AND LIQUIDUS TEMPERATURES

OF MAR-M 247 AND FIVE ALLOY

MODIFICATIONS,

Alloy Solldus Liquidus

Mar-M 247 1325sC (2417"F) 1383"C (2521eF)

t, AllOy 1 1345eC (2453"F) 1383"C (2521aF)

Alloy 2 1340sC (2444eF) 1410or (2570"F)

AllOy 3 1355ec (2471"F) 1390°C t2534 eF)

AllOy 4 1325ec (2417"F) 1380"C (2516°F)

AllOy S 1300"C (2372tF) 1390"C (2534aF)

DTA per forned at Earrop Laboratories in

Columbus, Ohio.

44

I

®
evaluated on Alloy No. i: 1260°C (2300°F)/2 hours and 1288°C

" (2350°F)/2 hours. The other alloys were all given the 1260°C

(2300°F)/2-hour solution treatments. These alloys were tested as

equiaxed cast ingots.

Metallographic characterization of the alloys and the heat

treatments performed are summarized in the following paragraphs.

Significant instabilities were not detected after either the 843 °

: or i093°C (1550 ° or 2000°F) exposures. Representative microstructures of as-cast, heat-treated and 843°C (1550°F) exposed

alloys are also provided.

5.2.1 Alloy 1

. _ Figure 15 depicts the alloy microstructure in the as-cast conbo

dition and after a 1260°C (2300°F)/2 hour heat treatment. Inspec-

: tion of the microstructure indicates that this temperature was too

.._'_ low to completely solution the gamma prime phase. In contrast, as

aP

shown in Figure 16 (top), a 2-hour exposure at 1288°C (2350°F) was ._

effective in solutioning the gamma prime phase. Small needles or

platelets, which were subsequently found to be _-tungsten, were also,

observed in the solutioned microstructure • 4_

. The microstructures at the bottom of Figure 16 show the _'.

effects of the 843°C (1550°F)/I176 hour exposure after the follow- t

ing heat-treatments:

o 982°C (1800°F)/5 hours (simulated coating cycle) + 871°C °

_ (1600°F)/20 hours (gamma prime precipitation)

O 1260°C (2300°F)/2 hours + 982°C (1800°F)/5 hours +

_ 871°C(1600°F)/20 hours ._

o 1288°C (2350°F)/2 hours + 982°C (1800°F)/5 hours + 871°C _

I (1600°F)/20hours

45 ,_

, ,l_ A
// ..

.._. _-

_,, _." , . _ • ,,.

(80X) "_o-,

Figure 15. (400X)

N_crostructures of Alloy Nodlficatlon No. 1.

UPPer Photos As-Cast, Lower Photos After

1260_C (2300_p) 2 Hour Solutlon Tceatment.

46
•..i+ (

tl t _'. tl _ , OR,GfN,+,_...... T "'"" .;.++.", _= : _.,

OF POOR QU,%LI+_y

* ++'_.:t'+. +.... _ . '.:. r _b. ' t'

,i" .+",,¢,P',,'".+; *'+ ,'' "' ° ++''+'/

.,+. ..+... ........ ,\ - +P

mMaim_+<_, - ,.., _:,'- +-++_. ,'_+._.+,. ... :r

..... ••r+,,=++ . + ++,%+++ _ ', +.+. ++. . ++'+_, +++,++.¢+ + +

., .,,. ,++. , +.. _;;_-/"+++_It_++,: '+ + +,+ ,<,. "+,

,=i__ _4,,,+. "}...++++

+++: + +'-+"++'++"'+" ' *

III _i+,i+++.,+,+

III ++

_, .: _ • . +_+m .1_-"+ ..+ _.-m+. . • + u+.. . . , +;.-+ •

"r'w"+"_,r,+' • " .+ ',; " ++ ' _t,.#+_'+ ': 'p ° : °

• "+"+ +_-w.. _ +"+ " " • .'_'++,,_+ " + _ .+.L+,"r+,'._ _I_, "' +_ " " " .'" "J+ "+- o;''"+'" " *+ ' " ++"" +

_., m_. -'+.++++ - ;++_+,._. ,' ..,,_++_,-++_'++ . .. _,+

._+ - *--,,,--' + + ,_@ , . +.,.,_ • , ," _s `+•+_ "' 7 _

+_ _,,,,mt__ ,++_,+, .+:-.+ .-++_+-+_,+' •:,_t_+ +++++ , . , +

/_.i _ "+ +"+"" ' '++'+++'_+ + ' + " "++'+ +'+..... '++ +

++.

•

t+.+_ _+' +iltW o• .4".-++.++" , _,_++-,,-,+' '.:,r+•,,

'+ ,,- +_++Y__. "M+"'J%+++Z+_P+,+'_4+_'++*

,

;_+++ + ">_+ " +S_I_,=

+ M" ,+ ++ " "; P_ ," ' "" C_+'++,+,_+'%'.'++'++b+_,,¢'",r.._

,f_.,.-"+ "/ _ "++_+"+,_; _" _':+_++. +.. +,+_:-+.+.

• • . ,,. _;== .., + , . -,, + + + _._,, ++,..+.

+',, _'., + , , _';_+ - , ,+ .= . ,,+',+_,++_,'

M+_.+._l_J_',+ _-'_T+_j_._,,-,._]-, .++ . ++-w ,. ,.",.%'Jgwa_,_I+, ............... -:. ' ' . :+

+ .;,-. ,,_+ +_ .,:,+ .+ +._,...,_.:' ,+.'+'=--"-+++. _+,;.+++" : ,, _.,<,+;:.+#+.-_+ - ,,.,,,, ._,

,_+ . + • , + .++- _ : +,,_ ;': + +_+ • ,,+. • , ,

+ .... +..... i

_+- " +. ._.... ,a.+ ,,_t_.._+++*.- " .-+:

i++" '+++ ' " _"++"++ '_' ++"_:'+"+" ?' ' +' ",. _ +: ,.,. ,., ., .+ . . _+L - + • .

.. • ' _ f." . . _,z,,_ ..re _ " ¢ " -

,,_+ _ -_-, . ., ++..--.+.+.,e+__; _ -.+.. . +

,..... , + • .,,,,+.,,_._ +++._-+.

'+_ +

.++p

, +

_++__

-+"

• +++",......'_

"+ "-

'++.,-

"

,

*

++. _ ++> "

+

.",

"_

*

" •

, +,

_

P...!.+_,'++.+ ++_++,,+ , '+_,_.:

';a

++_

•

,_+"+r+ +,

• . _+.+, .... , ._ ++,._ +,,,++.+ ++ "%" _ • +', P,,'+,+'++-+'+'+' _,++ ,r+-. ; ,.,,,,. ,,+e . -

,..,- .+ , ....

P78650-2 '-

Figure 16. Microstructures of Alloy Modification No. I.

Top: As-Cast + 1288°C (2350OF)/2 Hr '+

Solution Treatment (500X). ,'

Bottom: 100X After 843°C (1550°F)/I176 Hr

Exposure . i

Left: 982°C (1800°F)/5 Hours + 871°C [600OF)/ ",'

20 HES Prior to Exposure; ,:

Center: Same as Left but with 1260°C (2300°F)/ ,+

2 l+r Solution Treatment;

Right: S2amHeE aSsoluLetion ft buTrea t wtimen th t.1288°C (2350OF)/

47
Some grain bounda[j changes were observed for the last condition, but the matrix (present only in an SC blade) remained unaffected.

5.2.2 Alloy 2

Figure 17 shows the typical as-cast microstructure and modlfications of this microstructure resulting from a 1260°C (23000F)/2

hour heat treatment, and from a complete heat treatment. The

1260°C (23000F) temperature again appeared too low for effective

solution treatment. No effect of the coating and aging thermal

cycles was evident. Typical microstructures after thermal exposure

are shown in Figure 18. Some isolated needles or platelets were

_vident in the 1260°C (2300°F) treated material, which probably

formed prior to the 843°C (1550"F)/I176 hour exposure.

b

5.2.3 Allo z 3

Microstructures are shown in Figure 19 prior to exposure and "

in Figure 20 after exposure. The hafnium-modified eutectic gammaprime did not dissolve at 1260°C (2300°F). Some type of needle

phase formed around the "swirly" gamma-prime on the nonsolutioned

alloy, but the amount formed was not enough to cause concern.

5.2.4 Alloy 4

Structures prior to exposure are shown in Figure 21 and after

exposure in Figure 22. This relatively high Mo alloy modification

formed more of the needle phase during solution treatment than did

the other alloys. The 843°C (1550°F)/I176 hours exposure appeared

to have had no effect on the a?loy.

48

I

. ,, . - j ®
ORIGINAL PAG_

OF POOR QUALITY

•,' _' '_ ; )'_ " ,"_'lk" _6,,..[Kk_. .'_ - i

.'.' . _,, _,_ , ....t[' .-_)L,_:.'_

__:__,11_

.,,___',_. _ _' . ,. ,_ ' _:.

,.'.. , -'_-.... 4,', _' "_'T"..'i_'_,' "_

i,

Figure 17. Microstructures of Alloy Modification No. 2, _t

Top: As-Cas t. ;_,

Center: After 1260°C (2300°F)/2 Hr Solution '_

Treatment. :'!

Bottom: After 1260°C (2300°F)/2 Hrs + 982°C ,.

(1800_F)/5 Hrs + 871°C (1600_F)/ i:

20 Hours.

I

49 '_
OR|GINAL PAGE 18

OF POOR QUALITY

ot

.2

.L .. ,,

w

a

I

• . ,_.:,_,_,_n_'

' -_ ,_ _ _1_'='_

(80X) (400X)

Figure 18. Hicrostructu_es of Alloy Hodifica_.ion No, 2 after

843°C (1550°F)/1176 Hour Exposure.

Top: Prior Heat Treatment - 1260=C (2300°F)/2 Hrs +

982°C (1800°F)/5 Hrs + 871°C (1600 0F)/20 Hrs.

Bottom: Prior Heat Treatment - 982"C (1800°F)/5 Hrs

+ 871°C(1600"F)20 Hrs.

50

® m N
ORIGINALPA(_Erg

OF POOR QUALITY'
ORIGINAL PP,GE Ig

OF POOR QUALITY

• (75x)

\

• "'_ (375X)

• . JJ

t

Figure 20. Microstructures of Alloy Modification No. 3

after 843_C (1550_F)/I176 Hr. Exposure.

Top: Prior Heat Treatment - 1260°C (2200"F)/

2 Hrs + 982°C (1800OF)/5 Hrs + 871"C

(1600eF)/20 Hrs.

Center and Bottom: Prior Heat Treatment - 982_C

(1800°F)/5 Hrs. + 871eC

(1600°F)/20 Hrs.

52

I

® i
11

ORIGINAL P;.CL. ;.J

OF POOR QUALITY

;' ¢$

?.

,.

_. " • %Q "%,,,,. 175X)

I: • o

, |

,L 6 /

f

."

-,, 4-",.. 175X1

:'.

°" '_l

tl-

',t

e" p

-- )

,-., '-- ; (375X) ,_

;r 4

_" Figuge 21. Microstructures of Alloy Modification No. _..

_" Top: As-Cas t.

_'- Cente_ and Bottom: A_ter 1260°C (2300°F}/ '4

2 Hrs + 982°C (1800"F)/ _

-" ._ 5 Hrs + 8710C (1600°F)/

_ 20 HES.

53

II
_,L_._ ,* '"

";- _:_ _*,"'J_,"_!_ 1375X1

• . .. ,.

• : _ ''<' _ 1375X) ", _-.:,;._ ,=

t, P78650-,8

' Figure 22. Microstructures of Alloy Modification No. 4

: after 843_C (1550_F)/I176 Hr Exposure.

Top and Center: Prior Heat Treatment

982°C (1800°F)/5 Hrs

+ 871°C (1600°F)/20 Hrs.

Bottom: Prior Heat Treatment 1260°C

(2300°F)/2 Hrs + 982°C (1800°F)/

5 Hrs + 871°C (1600°F)/20 Hrs.

54
5.2.5 Alloy 5

This yttrium-bearing alloy formed a peculiar gray second phase

that is readily seen in the unetched photos of Figure 23. The

overall photos suggested that the yttrium was tied up as a lowmelting phase, some of which apparently incipiently melted at

1260°C (2300°F)/2 h_ur heat treatment. No indications of instability are evident in the photos of Figure 24.

5.3 SC Derivatives of Mar-M 247

Based on the "paper alloy" study and the furnace exposure

tests for microstructural stability, it was decided to procure

small master alloy heats and evaluate the five alloys in Table 13

• in exothermically cast SC form. This evaluation included estab- '_'

• lishing a heat treatment for each alloy and assessing their metal-

_':, lurgical and mechanical property characteristics. These aspects of

SC derivative alloy development from Mar-M 247 are reviewed in the

: following paragraphs.

|.

,>

5.3.1 Mast_or Alloys ,,

New master heats were procured from Cannon-Muskegon for this %

phase of the program. Aim and actual master alloy chemistries are _.

provided in Table 15. Alloy 5 was produced using a late addition <:

of yttrium to the Alloy 1 master heat during subsequent melting at .-:

Jetshapes. The procedures to produce the master alloys are

reviewed in another publication. 13

P _

5.3.2 SC Casting Campaigns ':

' ,¢_

i AS discussed in Section III, the initial molds of the ',:

Mar-M 247 derivative SC alloys were contaminated with zirconium '.:_

:' during melting in a zirconia crucible. Thus, these SC alloys could 4,

_ not be adequately solution heat-treated and resulting strengths V

were below DS Mar-M 247. '_

, it

55

i 
Figure 23. Microstructures of Alloy Modification No. 5.

Top: As-Cas t.

Center and Bottom: 1260°C (2300°F)/2 Hr

Solut ion Trsatment.

56

I

_,_" ...... ,
++

(,

!

ORIGINAL PAGE +.8 L

. OF POOR QUALITY

• : "": . L. "'+

g : +': : ;""+'+ " +:" IOOX

+-+

d : : .

l"

. .'

..... ; .., +- _ ,_. +'_. .-.'--,dr,w,,v_ _j 4+ _- .+ <

w, :. ...;. . . . . ,. .., : -+._+__. .

. 2 3,1 -''+

,,. . + '......,,.+; _

ii . ,f , , • . ¢ +

.- "+_ . ,+'+ " " ' "_ . ", P. _'_'+."'_+ ,m",++ ,:.

+ , +. , .+++v+",e

-: + : ' " °,b",,........ ,+t+

-

_..+++ .++;

;' IOOX :,. .I

,+4

! .," 4_ . '+

-. %+:-,

Figure 24. Microstructures of Alloy Modification No. 5 after '

843°C (1550°F)/I176 Hr Exposure. _

Top: Prior Heat Treatment - 982°C (1800°F)/ 5 Hrs ,t

+ 871°C (1600°F1/20 Hrs. L

Bottom: Prior Heat Treatment - 1260°C (2300OF)/ _! i

2 Hrs + 982°C (1800°F)/5 Hrs + 871°C Y (1600°F)/20 Hrs.

57 ':
ORIGINAL PAGE IS

OF POOR QUALITY

TABLE 15. COMPOSITIONS OF MAR-M 247 ALLOY MODIFICATIONS MELTED

BY CANNON-MUSKEGON.

Alloy Composition (Weight Percent)

_ Alloy 1 Alloy 2 Alloy 3 Alloy 4

q.

_' !VF-31 VF-154 VF-100 VF-99 VF-98

Element Aim &ctual Actual Alm Actual Alm Actual Aim Actual

.........

!MO 0.8 0.8 0.0 0.8 0.8 0.8 0.8 2.0 2.1 f

lW 10.5 L0.2 10.4 10.2 10.0 9.4 9.2 10.1 9.9

ITa 3.3 3.4 3.4 3.4 3.5 3.2 3.3 3.4 3.5

i TI 1.2 I.i I.I 1.1 1.0 1.0 0.93 1.1 1.0

_; Cr 9.0 9.0 9.0 9.0 9.3 8.5 8.2 8.9 9.0

!. Co .......... 5.0 5.1 ....

q A1 5.75 5.5 5.8 5.75 5.8 5.45 5.54 5.65 5.74

. C -- 0.018 0.0034 0 0.0032 0 0.0018 0 0.0022

Hf ...... 0.6 0.6 0.6 0.6 ....

-b Zr LAP** -- 0.01 LAP -- LAP ....

Oxygen* -- 4 .... 5 -- 6 -- 6

Nitrogen* -- 6 .... 3 -- 14 -- 4

-;- Ni Bal Bal Bal Bal Bal Bal Bal Sal Bal

**Low as possible _-

":" Actual compositions provided by Cannon-Muskegon

-_ This above problem was solved by melting the alloys in a magnesia crucible instead of a zirconia crucible. A second casting

• campaign was initiated to evaluate SC alloy modifications 1 to 4.

_ (Alloy 5 was dropped from further consideration at this time. Both

the zirconium contamination and the yttrium addition were concentrated in the incipient melting zones.) As in the previous casting }4

campaigns, the DS blade configuration from MATE Project 1 was used

for the casting evaluation. Single-crystal TFE731 high-pressure

turbine blades were cast with Mar-M 247 and modified Alloys i, 2,

3, and 4. A magnesia crucible was used to melt the alloys at a

i nominal temperature of 1593"C (2900°F). Ten molds were cast, each

containing twelve SC blades. The results of visual examination of

the product of these molds are shown in Table 16; 64 percent of the

blades, which filled with metal, were single-crystal with the

desired [001] grain orientation within 15 degrees of the blade's

longitudinal axis. SC Alloys 1 and 3 appeared to have the best

castability based on this limited data.

; 58

® q_ ..,°

,¢
'i". TABLE 16. VISUAL INSPECTION RESULTS FOR SC TFE731 TURBINE

BLADES OF MAR-M 247 AND MODIFIED SUPERALLOYS.*

Blades Filled** Single-Crystal

Alloy (Mold No.) With Metal AIrfoil ard Root***

. . , . •

Mar-M 247 (4671 8 2

MaE-M 247 (468) i0 7

;_ Alloy 1 (4581 12 8

Alloy i 14811 12 i0

Alloy 2 (465) 5 2

J

:; Alloy 2 (482) 12 6

- Alloy 3 1464) i0 7 _

Alloy 3 (483) 12 9 .

Alloy 4 (466) 8 6 r

Alloy 4 1484) 8 5

:. 97 62 I

• _

5_. *Cast with MgO crucible.

_,. **Each mold contained 12 blade cavities.

_2 ***[001] grain orientation within 15 ° of blade longitudinal axls.

_ A comparison of master melt and SC blade chemistries is

' . provided in Table 17. These results indicated that magnesium _-_

_ pickup ranged from 2 to 12 ppm; magnesium additions at this level

were considered innocuous. Changing from a zirconia to a magnesia

• crucible to melt Alloys i through 4 reduced the zirconium in the SC

,:- blades by a factor of about 40. _

_- 5.3.3 Heat Treatment Study :_

e ,

_" The solution heat-treatment ranges were assessed by furnace \:

:4

' tP •

_, exposure followed by metallographic examination. Results of this '_

;_ evaluation are summarized in Table 18. As-cast and solutioned i

'* microstructures of Alloys I through 4 are presented in Figures 25 ,._

to 28, respectively. Based on these results, Alloys 1 and 4 were _

solutioned at 1316"C (2400"F) for 2 hours and 1324°C (2415"F) for _

2 hours. The hafnlum modified Alloys 2 and 3 exhibited incipient '%

_ melting at lower temperatures and were solutioned at 1288°C _

i1_ ''_''_'''_'_ "-" -,
!

!

• ORIG;NAL PAGE IS

; OF POOR QUALITY

i

TABLE 17. MASTER MELT AND BLADE CHEMISTRIES OF MAR-M 247 AND MODIFIED MAR-M 247

SINGLE-CRYSTAL ALLOYS MELTED WITH A MAGNESIA CRUCIBLE.

I ! Actual Alloy Composition (Weight Percent)

i

!

Mar-M 247 Alloy 1 Alloy 2 _ Alloy 3 Alloy 4 l

f_ Element Range Blade Mast Melter Blade MasMeltter Blade MMelt aster Blade MaMseltetr Blade

Mo 0.50-0.80 0.7 0.8 0,8 0.8 0.8 0.8 0.8 2.1 2.1

w 9.50-10.50 9.8 10.4 10.7 10.0 10.0 9.2 9.4 9.9 10.2

Ta 2.80-3.80 3.0 3.4 3.5 3.5 3.6 3.3 3.3 3.5 3.6

Ti 0.90-1.20 1.0 1.i I.i 1.0 1.0 0.93 0.9 1.0 1.0

,!" Cr 8.00-8.80 8.0 9.0 7,5 9.3 8.8 8.2 .7 9.0 8.6

Co 9.00-11.00 10.1 ........ 5.1 5.1 ....

A1 5.30-5.70 5.4 5.8 5.6 5.8 5.7 5.54 5.6 5.74 5.6

> C 0.13-0.17 0.13 0.003 0.0012 0,003 0.0019 0.002 0.0016 0.002 0.001_

Hf 1.20-1.60 i.i -- <0.01 0.6 0.6 0.6 0.6 ....

Zr 0.03-0.08 0.03 0.01 0,002 0.014 0.017 0.005 0.009 0,001 0.002 i

B 0.01-0.02 0.016 ................ I

m

- Mg -- 0.0002 -- 0.0005 -- 0.0012 -- 0.0006 -- 0.0004

Ni Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal "_

D

NOTE: Master melt and blade chemistries were determined by Cannon-Muskegon.

L

TABLE 18. SOLUTIONING HEAT-TRF-%TMENT CHARACTERISTICS OF MODIFIED MAR-M 247

SC ALLOY TFE731 HP TURBINE BLADES. a

Solution SC Alloy Modification Number ...........

| Treatment ........................................

___°F)/Hours I 2 3 4

• J1288 (2350)/4 Partial solutioning --- Partial solutionlng ___

partial solutlonlng Partial solutioning, .................... Partial solutionlng, -85t solutlonlng

_, i incipient melting incipient melting

1316 (2400)/2 -85t solutioning 80t solutionlng, 85% solutioning, -951 solutloning

incipient melting incipient melting

1316 (2400)/2

1329 (+2425)/2 >90t solutloning ...... Complete solutionin

-, .....................................

i', 1343 (2450]/2 Incipient melting ...... Incipient melting

...........................

1288-(23"50)/2 --- Incipient melting Incipient melting ---

g

, _ 1316 (+2400)/2 ..................................................................

_' 1288 (2350)/2 _ 75t solutioning, 85% solutioning .-.

1302 (+2375)/2i "-" incipient melting ...................................... incipient melting

• SO alloys melte(] in megnesia crucible a.d cast at fletshapes, Inc.

........................................

°.

L 60

. . -_ . m,__-_ "

m

_.m , • ...... , . v._ .... --,,w--, ,,me, , , _,, ,• ,, ,..... ,

i R, ............... _, - _-..•..:...e_ _ ,. ,_;,
ORIGiNA:. " ......

OF POOR " .... "'"

il

' AS CAST

o

b

|

2-

- ,_, :.->. -

, .. ... 1288°C (2350°F1

: '*_ . '_::_ FOR 4 HOURS "

J IlL

• ,. 12400OF1 _-.

• ,-, _'., ' ." FOR 2 HOURS

, AND

• 1329Oc (2425OF)

.. , /_ . FOR2HOURS '-_,

'9

- " P786EO-11

Figure 25. Microstructures of Alloy 1 Turblne Bla_es

:_ in As-Cast and Solution Heat-Treate_ _"

Condi tions. _i_

_ ,
f

=J

!

ORIGINALPAGE [g

OF POORQUALITY

_ AS CAST

¢.

4.

,o

r

.

.. _i'! _ _

- ,,. ,.-,' 1302°C 12375OF1FOR

:..,. _ _'%"_i;'_ 2 HOURS

I 4 '

....._ __

P78650-12

Figure 26. Hicrostructures of Alloy 2 Turbine Blades

in As-Cast and Solution Heat-Treated

Conditions.

62

I

®
ORIGINAL PAGE IS

OF POOR QUALITY

,'% ,

q ,/ _',*

" ' AS CAST

i

,: " 2 ' ", , 1288oc 12350OF1 FOR

._ '* 4 HOURS

.......... I,

"t

" <. 1288°C (2350OF) FOR 2

; HOURS

1302°C (2375°F) FOR 2

HOURS ' ',

1=__1 "

P78650-13 "

!,

Figure 27. Microstructures of Alloy 3 Turbine Blades ¢

in As-Cast and Solution Heat-Treated ,_

Cond it ions. :i

63

, !

.......®i
ORIGINAL PAGE IS

OF POOR QUAL|3"_

i,

;:- AS CAST

li"

I,

i

I

*_o- , " 1316°C (2400°F1 FOR 2 HOURS

¢

#

1316°C 12400°F) FOR 2 HOURS AND

1329°C (2425°F) FOR 2 HOURS

P78650-14

Figure 28. Hie_ostructures of Alloy 4 Turbine Blades

in As-Cast and Solution He_t-Treated

Conditions.

64

i
E

i

,i

ORIGINAL F_ IS

OF POOR QUALITY

.. (2350°F) for 2 hours and 1296°C (2365°F) for 4 hours. EDX analysis

indicated that the lower incipient melting temperatures of Alloys 2

and 3 were associated with a hafnium rich eutectic; magnesium was

not detected in incipient melting zones.

• For SC Alloy i, the solutioning temperature was significantly

: above the 1288°C (2350°F) used earlier for the equiaxed casting

(Figure 12). This result Js attributed to the significantly larger

; dendrite arm spacing of the SC alloy relative to the equiaxed

o alloy.

Solutioned [1246°C (2275°F) for 2 hours in vacuum] microstructures of SC Mar-M 247 blades, which were cast with the alloy melted

•" in magnesia and zirconia crucibles, are compared in Figure _:.

Both microstructures were similar; solutioning of the 7' was incomI.

plete and a small amount of incipient melting was present. SC

w -

Mar-M 247 was not adversely affected by the zirconia crucibles.

At the conclusion of the vacuum solution heat treatment, the t

l

blades were fan-Argon quenched. Insptction of solution heattreatment records indicated that approximately 1.7 minutes were

required during fan-Argon quenching to reduce the temperature from

1324°C (2415=F) to 816°C (1500°F). Im --

4

I

%

: MAGNESIA CRUCIBLE ZlRCONIA ,_RUCICLE P_0-1§ .z

USED TO MELT ALLOY USED TO PELT ALLOY

Figure 29. Solutioned - 1246=C (2275°F) for 2 Hrs in Vacuum -

- _ Microstructures of Mar-M 247 SC Turbine Blades. ..,
Fan-Argon quenching was found to be a critical part of the

solution heat treatment when Jt inadvertently was not activated

until a batch of specimens had cooled to I038°C (1900°F); in this

case, 4.4 minutes were required to cool through the 1324 ° to 816°C

(2415 ° to 1500°F) range. The relatively slow cocl_ng rate from the

solutioning temperature resulted in a coarser (lower-strength)

gamma-prime precipitate morphology as shown in Figure 30.

5.3.4 Property Characterization

The SC alloy dezivatives of Mar-M 247 were screened for

tensile and stress-rupture strengths, oxidation resistance, and

density in order to facilitate selection of two promising alloys _

for design data generation and subsequent engine testing. This

effort provided data for both properly processed SC alloys and

deviate material (SC alloys contaminated with Zr from the crucible

and SC components that were improperly heat treated). Subsequent

paragraphs review the properties of the properly processed material

i

followed by a discussion of the effects of improper pro ssing on

SC alloy properties.

5.3.5 Tensile Tests

Longitudinal tensile test data was obtained for SC Alloys 1

through 4 at 760°C (1400°F), a temperature typical of blade attachments. As indicated in Table 19, SC Alloys i, 2, and 4 exhibited

average yield and ultimate strengths of 1235 (179) and 1366 MPa

(198 ksi), respectively. For reference, average yield and ultimate

strength levels for DS Mar-M 247 were 9;3 (141) and 1235 MPa

(179 ksi), respectively. SC Alloy 3 tensile properties were similar to those of DS and SC of Mar-M 247 at 760°C (1400°F).

!

66

V 4. -_ A
|'

,

Yt

,| ' :I

I

,:_

t, ' 67
!

!

,' ORIGINAL PAGE !_

OF POOR QUALITY

TABLE 19. 760"C (1400°F) TENSILE TT_ST RESULTS ON 1.8 Mel (0.070 INCH) DIAMETER

- LONGITUDINAL TEST SPECIMENS MACHINED FROM SINGLE-CRYSTAL TURBINE

BLADES OF ALLOYS 1 THROUGH 4.

UTS 0.2t Y.S.

_pecimen EL RA

No. Alloy Orientation MPa (ksl) MPa (ksl) (%) (%)

+

, 481-3L 1 Longitudinal 1379 (200) 1289 (187) 3.7 11.8

481-7L 1 Longitudinal 1407 (204) 1338 (194) 6.0 14.2

482-7L 2 Longitudinal 1310 (190) 1234 (179) 4.8 15.3

482-10L 2 Longitudinal 1386 (201} 1269 (184) 8.5 14.0

I

:- 483-7L 3 Longitudinal 1200 (174) 930 (136} 9.5 14.4

483-12L 3 Longitudinal 1179 (171) 958 (139) 7.7 15.2

• 484-5L-2 4 Longitudinal 1413 (205) 1165 (169) 3.8 13.3

" ,

l 484-6L-2 4 Longitudinal 1289 (187) 1103 (160) 3.4 9.7 _:

y-+

Heat Treatment:

a. Alloys 1 and 4 - 1316"C (2400"F} for 2 hours, plus 1324"C (2415"F) for

2 hours, plus 982°C (1800*F) for 5 hours, plus 871"C (1600"F) for 20 hours.

b. Alloys 2 and 3 - 1288"C (2350"F) for 2 hours plus 1296"C (2365"F) for

_ 4 hours, plus 982°C (1800"F) for 5 hours, plus 871 *C (1600"F) for 20 hours.

NOTE:

Alloys were melted in a magnesia crucible prior to casting.

5.3.6 Stress-Rupture Tests

i Stress-rupture data were obtained for SC Alloys 1 through 4 +

,. and SC Mar-M 247 at 760 ° (1400°), 982 ° (1800o), and i093"C (2000°F);

test results are provided in Tables 20 to 22 for the respective

:' temper res.

Examination of the data indicated that all of the modified SC

_, alloys were significantly stronger than SC Mar-M 247. SC Alloys i,

,. 2, and 4 exhibited excellent rupture lives under these test conOi- +

tions, while SC Alloy 3 was only slightly weaker at all temperatures. Based on the Larson-Miller extrapolation, at a 103 MPa (15 +

ksi) stress level, SC Alloy 1 had 58 ° and _"C (iU5 ° and 5°F) advan- i

tages in temperature capability (for a given rupture life) relative !

.- to SC Mar-M 247 and SC Alloy 4, respectively. At the 207 MPa (30

68

i
' ORIGINAL P__V"_._ r_

OF POOR QUALITY

TABLE 20. 760°C (1400"F) STRESS-RUPTURE RESULTS ON 1.8 MM (0.070 INCH) DIAMETER

SPECIMENS MACHINED FROM SINGLE-CRYSTAL TURBINE BLADES.*

|

.................................... _ ............................

RUPTURE

SPECIMEN TEMPERATURE STRESS LIFE ELONCATION R of A

NUMBER ORIENTATION ALLOY "C (°F) MPa (ksi) (HOURS) (%) (%)

.................................

_. 467-3L LONG. Mar-M 247 760 (1400) 724 (I05) 213.4 6.8 15.2

468-6L LONG. Mar-M 247 724 (105) 276.6 7.9 16.8

,_ 467-3T TRANS. Mar-M 247 724 (I05) 338.7 I0.0 20,9

v

j 468-6T TRANS. Mar-M 247 724 (105) 32.7 9.6 16.2

_ 1 481-5L LONG. 1 758 (ii0) + 793 (i15) 790.2"" 6.9 13.2

481-12L LONG. 1 758 (110) + 793 (i15) 725.2** 7.1 15.4

- 481-5T TRANS. I 758 (110) + 793 (115) 551.8"* 10.8 23.2

• 481-12T TRAN3. 1 758 (II0) 308.6 3.8 I0.0

_ ", 465-4L LONG. 2 758 (II0) + 793 (I15) 618,5"* 6.8 19.8

|' 482-4L LONG. 2 _58 (Ii0) + 793 (115) 622,9** 5.9 14.7 -'

_ 46_-4T TRANS. 2 758 (II0) 405.5 5.3 12.5

=. 482-4T TRAt_S. 2 758 (II0) 393.6 2.9 11.0

•. !4RS-[L LONG. 3 758 (110) 317.8 15.4 23.7 I

!483~2L LONG. 3 758 (110) 291.8 16.9 25.3 I

!484-5L LONG. 4 758 (110) + 793 (I15) 894.2"* 8.0 15.3 ,:

J484-6L LONG. 4 758 (i10) + 793 (i15) 773.2** 6.6 15.2 _'

4- / ............................

. . *Heat treatments:

a. Mar-M 247 - 1246"C (2275"F) for 2 hours wLth Argon Quench, plus 9820C (18000F) for 5 hours, plus

_* , 8710C (1600"F) for 20 hours.

; b. Alloys i and 4 - 1316"C (24000F) for 2 hours, plus 1324"C (24150F) for 2 hours with Argon Quench,

plus 9820C (1800"F) for 5 hours, plus 871°C (1600"F) for 20 hours.

-: c. Alloys 2 and 3 - 1288"C (2350°F) for 2 hours plus 1296°C (2365°P) for 4 hours with Argon Quench,

. plus 982"C (1800"F) for 5 hours, plus 871°C (1600°F) for 20 hours, i .

**Specimens were uploaded from 758 (ii0) to 793 MPa (115 ksi) after 500 hours. .@ . -.

%1

- _ ksi) stress level, SC Alloy 4 had 18 ° and 6°C (32° and II°F) advan- :

tages relative to SC Mar-M 247 and SC Alloy i, respectively. SC -

= Alloy 2 rupture lives were intermediate to those of SC Alloys 1 and _

: 4 At 760°C (1400°F), SC Alloy 4 also exhibited the longest .

rupture lives. &

?

Q

_ Transverse [I00] stress-rupture data were obtained for SC

Alloys 1, 2, and SC MJr-M 247. Generally, the rupture lives of the "

transverse [i001 material were approximately equivalent to the •

_% longitudinal [0011 material.

Ductilities of all the SC alloys tested were jatisfactory for

turbine blade applications. 'P_ I

t

e=

69
J.

I-

|

T

ORIGINAL P,_,_ i_

OF POOR QUALITY

TABLE 21. 982°C (1800"F) STRESS-RUPTURE RESULTS ON l.S MM (0.070 INCH} DIAMETER

_PECIMENS MACHINED FROM SINGLE-CRYSTAL TURBINE BLADES.*

RUPTURE

SPECINEN TEMPERATURE STRESS LIFE !ELONGATION R of A

NUMBER ORIENTATION ALLOY "C ("F) NPa (ksL) (HOURS) (%) (t)

_0 .,

468-5L LONG. Mar-M 247 982 (1800} 207 (30} 123.5 23.6 45.7

468-7L LONG. Nar-M 247 120.5 27.4 44.5

468-5T T°ANS. Nar-M 247 115.4 36.4 53.9

_. 468-7T . .ANS. Mar-M 247 116.4 27.8 46.0

_ 458-3L LONG. 1 i 189.3 24.3 42.6

458-4L LONG. 1 198.1 15.7 41.3

i 481-3T TRANS. 1 174.4 15.6 26.3

481-7T TRANS. 1 189.5 26.0 34.2

_! 482-6L LONG. 2 216.8 21.1 36.8 |

: 482-3L LONG. 2 217.1 19 •2 43.6

._ 482-6T TRANS. 2 164.9 8.5 26.7

.:. 482-3T TRANS. 2 258.4 8.9 33.3

• 483-3L LONG. 3 180.2 22.5 28.9

483-8L LONG. 3 161.3 10.2 28.3

484-IL LONG. 4 251.4 23.7 30.4

484-3L LONG. 4 1 I 249.7 24.8 37.2 ,_ P

_:

_ TABLE 22. 1083°C (2000°F) STRESS-RUPTURE RESULTS ON 1.8 NM (0.070 INCH} DIAMETER

• SPECIMENS MACHINED FROM SINGLE-CRYSTAL TURBINE BLADES. _

Rupture

. Specimen Temperature Stress Life Elongation R of A _

_ Number Orientation Alloy "C ("F) MPa (ksi} (hours) (%) (_} ." _

e

x

_ 468-4L Long. Mar-M 247 1093 (2000) I03 (15) 72.4 22.4 66.5

468-9L Long. Mar-M 247 39.0 27.2 49.9

= 458-6L Long. 1 501.8 11.2 34.5

_ 458-IIL Long. 1 453.7 10.8 33.7

; 482-1L Long. 2 456.8 _ 7 40.5

L,

_- 482-2L Long. 2 458.4 5 28.2

483-9L Long. 3 247.8 ll.a 20.0

i

483-IOL Long. 3 I 199.1 16.9 31.0

i

484-2L Long. 4 399.9 II.0 30.7

I

• 484-4L Long. 4 i 457.2 16.3 27.1 __ i

•Heat treatments:

a. Mar-N 247 - 124C"C (2275°F) for 2 hours wlth Argon Quench, plus 982"C (I800"F) for

5 hours, plus 871eC (1600"F) for 20 hours.

b. Alloys 1 and 4 - 1316eC (2400eF) for 2 hours, plus 1324eC (2415"F) for 2 hours with

Argon 0uench, plus 982eC (1800*F) for 5 hours, plus 871"C (1600°F) for 20 hours.

c. Alloys 2 and 3 - 1288"C (2350eF) for 2 hours plus 1296eC (2365"F) for 4 hours with

. Argon 0uench, plus 982°C (1800oF) for 5 hours, plus 871°C (1600"F) for 20 hours.

' 70

I

t..

i
Typical microstructures of SC Alloys 1 to 4 in the pretest

condition and after evaluation in 760 ° , 982 ° , and I093°C (1400 ° ,

1800 °, and 2000°F) stress-rupture tests are provided in Figures 31

to 38. Inspection of these figures indicated that, with the exception of Alloy 2, only traces (if any) of the eutectic _' remained.

Only minor amounts of incipient melting were observed in the microstructures of SC Alloys 2 and 3.

In addition, the dendritic regions of SC Alloys i, 2, and 4

exhibited numerous small particulate and acicular phases that

energy dispersive X-ray (EDX) analysis indicated to be rich in i

tungsten. EDX analysis showed that the acicular particles also

containel some nickel. ,_

Results from Dr. J. F. Radavich (Micromet Laboratories, W.

Lafayette, Indiana) of the microstructural analysis of stress-

" rupture tested MFB bars of SC Alloys 1 and 4 indicated that the ?

,. tungsten-rich particulate and acicular phases were u-tungsten and _

mu phases, respectively This was consistent with analysis of _

blade alloy chemistries by Cannon-Muskegon that indicated the

carbon contents of the blades were less than 20 ppm. At this low

carbon level, the probability of an M6C tungsten-carbide formation

was extremely low.

Examination of the post-test microstructures of the stressrupture specimens indicated that secondary creep cracks generally

initiated at small pores in the interdendritic regions. [The

interdendritic regions were about 20 percent lower in tungsten

t ':

%

-!

r 71
¢

I

,)

t

ORIGINAL PAGE I_

OF POOR QUALITY

," q ,t _ -. ; _,- %

"

• ,_:_,::_'_ .., ,. , ,*s_=_. , .

E

_(' P78644-1 P78644-2

. a. PRETEST MICROSTRUCTURE

... .. _,:_ ,_._ ,% *t..-_'.c,_-_._._s.c_a_ •

,_,--; .... _,;-,_ ;.', ., ,.

: . ._..- _.., ,°. , :,.,:. . .....,;,.,, ,, ..... _, .'_;_ _._:_9:)t,_,,

_ :i :_' ...... ,_, ._,g"

t

if

.. ", '- "%_-,. _" .',. ,"._,i_ - _,'_.' ,_, • ..... "-_'_,#,_,_I

', ' _ " "_, hnnuml • % ' ..;c_>'./_, '., .,_,_" :_ ;_'-.._...,. ,.,= ....,

P78644-3 P78644=4

b. 760°C (1400°F) - 500 HOURS FOR 758 MPa (110 KSI)

PLUS 290.2 HOURS FOR 793 MPa (115 KSl)

". Figure 31. Pretest and Stress-Rupture Tested Microstructures

- of SC Alloy i, Part i.

72
.!

" OF POOR QUALITY

• ' /........ I __• _" _''* _ " ";**;V "" "" "_l_kA_.'_"%."" "_ '

.," . • _'/- ... '.:',: _.,,,

.. . _.. "o ,;* -. . + ,>_;..... ; ' t f ",_ %._. ,

.,._: ,_.., . • ;_,_.'_ -..

• t" ,: • • - "

•' '- -:_,: ",t-- " .-" ,. ". ._,.""

_. r.,

_" P78643-1 P78643-2

_:'_ c, 982°C (1800°F)--189.3 HOURS AT 207 MPa (30 KSI)

-

.... , :..II_ " ,-.%___,_ _._• - , , .-.

.. , • . _, .,.... . ._.,._. . .,_ _.__--- .; _ .,._-'- _ • . ,_ ...

"" " •., ( _ .',,. "%: . , . _" "_ ./ 0 ._ t

..... . j. ,-- .., ,_ ,

" ,.... - _'..,,., _:,.;. .'. ...._'_,.." .' '..... , " ...'..:-,-_ _ t :_.;" -

_ ,.'.;, ,., _ "I.C,-_,; _. • .. . . ._ $ ",, . .. _..._-_;.. _ ..: _-_ ,

._ . _. _.." • -_': ,.._r_ : " • . - : • ..... _. . ._ - - _ ..

.- . "', .,'_', _ _, ,. . _ , _, ... _- . _. , ." ._.

"' _ " /'_'_ "'¢" _'#_' _':"'_'" ; _ "' "..... " " "--a'.""'""" _' '_X" " o'-

• "_:. ,., .,, _:+. • , ;_._ . . .._:. .,,, _,,,, :;._ . ,. .. ...: ,S._::_ .; ,, .. :_'_?_ ._ ,: ..._.... . ..... _ . . ..........

'" """ _, " . h.__. "-'._'_"-_. , ' :_'", "", ::-_;:... .-' ,_ _-_ v. _'.__" ;-__'- _. .," . _--.-_

,'. _;,_;, .,.]i_. _.Z'" " _, ,._'_.,_ ', :Z.'_-:..,.'_.'G"__. -.-" _:_._.G_',_"_,..'- '.- - l,_,,'P,,q :

P78643-.3 P78643-4

d. 1093°C 12000°F)-501,8 HOURS AT 103 MPa 115KSl) :'

!" Figure 32 Pretest and Stress-Rupture Tested Microstructures ,I

! " _

-_r 73 1
oKIGINAL pAGE IS

OF poOR QUALITY


ORIGINAL PAGE P3

OF POOR QUALITY
I

1"|,)

ORIGINALPA__

OF POOR QUALITY
.

4

,'ORIGINALP_G_13 ' OF POOR QUALITY
!

r_

!

than the dendrite cores and wer_ free of tungsten-rich secondary

phases.]

5.3.7 Oxidation Evaluation

Burner-rig cyclic-oxidation tests were conducted at I093°C

(2000°F) for i00 hours. Pro- and post-test measurements of the

;:" specimens indicated t _t SC Mar-M 247 and Alloys 1 through 5 all

exhibited metal losses less than or equal to 25 _m (0 001 inch).

•

t: Metallographic examination indicated tnat, in addition to the

_ material loss, a y' depletion zone was present in the alloy adjaL. cent to the surface; the thickness cf ithe alloy depletion zone

varied from about 25 to 64 _m (0.001 to 0.00Z5 inch). _his test

b confirmed that the SC alloys would require protective coatings for ,_

, use in advanced commercial engines. (Protective coatings are oval- _'

C uated Jn the next section).

5.3.8 Density

L

Densities of SC turbine blades ca_t from Mar-M 247 and Alloys

: 1 through 4 as determined by the water immersion method are pre-

• sented in Table 23. These densitie_ were in excellent agreement

with predicted densities from the paper alloy study.

• TABLE 23. DENSITIES OF SINGLE-C[:_STAL TFE731

:, TURBINE BLADE CASTINGS. j

• Density

_% ! 3

Alloy gm/cc i (ib/in.)

I

l-

' Mar-M 247 8.62 (0.311) I

Modified Alloy 1 8.62 (0.311) !

Modified Alloy 2 8.62 (0.311)

Modified Allcy 3 8.62 (0.311)

Modified Alloy 4 8.63 (0.312)

80

I
5.3.9 improper Processing OF POOR QUALI_

AS noted pz:eviously, mechanical property dat:a was obtained in

two instances for SC alloys that were improperly processed. The

,+ first instance involved zirconium contamination (from a zirconia

,_- crucible), while the second was associated with slow cooling from

_" : the solutioning temperature. Stress-rupture test results at 1393°C

i (2000°F) and 103 MPa (15 ksi) are provided in Figure 39. These

"I: results illustrate the magnitude o_ life reduction associated _-i_h

!. the improperly processed material.

"_o

Test data for zirconium _ontaminated SC Alloys 1 through 5 are

:-+:

+ provided in Tables 24 to 2_. r _pecti,_ely, which show that not one

." ._j

_'. of the alloys realized its rupture life potential. This result is

i:':" predominantly attr!' teC to use of a nonoptimum solution heat

_5

-21. treatment temperature, 1260°C (2300_F), to avoid substantial

,: incipient melting. ;, %

_ %_ ,..

":+i STRESS-RUPTURE CAPABILITY{JFSC ALLOY1 _: --

:" bOO] '

i : 1093°C12900°FlAND,03 ',5 KSll "_ " . 4nO_ :_ !

_-" : 300-

e _

:_ m"::ZOO- IMPROPERLYPI_OCESSED , i

•_- ,.- SC ALLOY1 .

: = / ",, ,. I

'°°1 / :" '

,+

•t OSCMAR-M I-I 24?_SCN._SAI3I00 ZiRCONIA r-1 DELAYED '::,_'°

CRbCIBLE ARGON , _:$

_UENCH :_.

Figure 39. Effect of Improper Processing on Stress-Rupture :'_

--_ Capability of SC Alloy i.

#'k

4

l
ORIGINAL P:_,GE'lg

OF POOR QUALITY

TABLE 24. STRESS-RUPTURE TEST RESULTS ON 1.8 _91 (0.070 INCH)

DIAMETER LONGITUDINAL TEST SPECIMENS MACHINED FROM

SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED*

ALLOY l+Zr.

F

Rupture

; Specimen Temperature Stress Life Elongation R of A

_ Number eC (*F) MPa (ks/) (Hours) (%) (%)

• 426-1L 982 (1800) 207 (30) 71.8 13.9 24.8

_ 426-3L 982 (1800) 207 (30) 79.0 22.6 45.6

I 426-2L 1038 (1900) 131 (19) 121.2 17.3 34.7

426-5L 1038 (1900) 131 (19) 113.8 17.5 39.2 i

426-9L 1038 (1900} 131 (19) 129.6 22.2 36.1

426-4L 1093 (2000) 103 (15) 41.2 29.5 58.2 ._

P

426-75 1093 (2000) 103 (15) 43.4 30.0 49.8

_ *1260eC (2300eF) for 2 hours with Argon 0uenoh, plus 982eC (1800°F)

, for 5 hours, plus 871"C (1600eF) for 20 hours.

?

_ TABLE 25. STRESS-RUPTURE TEST RESULTS ON 1.8 MM (0.070 INCH)

DIAMETER LONGITUDINAL TEST SPECIMENS MACHINED FROM '

SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED*

ALLOY 2+Zr.

Rupture

: Spech,en Temperature Stress Life Elongation R of A '

_ Number eC (eF) MPa (ksl) (Hours) (%) (%)

' - . ,

433-il 982 (1800) 207 (30) 7_.8 44.6 54.1

b=

433-5L 982 (1800) 207 (30) 72.2 28.1 44.8

: 433-6L 1038 (1900) 131 (19) 109.7 30.5 53.7

433-8L 1038 (1900) 131 (19) i06.1 29.2 52.6

433-2L 1093 (2000) 103 (i_ 25.2 32.8 58.8

_. 433-7L 1093 (2000) 103 (15) 38.3 29.7 59.1 ,

_c '

*1260eC (2300"F) for 2 hours with Argo.1 0uench, plus 982e_ (1800"F)

:_ for 5 hours, plus 871"C (1600"F) for 20 hours.

82

®
!

ORIGINAL P;_C-'__"

OF POORQUALITY

TABLE 26. STRESS-RUPTURE TEST RESULTS ON 1.8 HM (0.070 INCH)

DIAMETER LONGITUDINAL TEST SPECIMENS MACHINED FROM

SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED*

ALLOY 3+Zr.

Rupture

- ' Specimen Temperature Stress Life Elongation R of A

Number "C ("F) MPa (ksi) (Hours) (t) (%)

428-9L 760 (1400) 724 (1051 132.1 14.9 15.4

447-8L 760 (14001 724 (1051 141.4 12.8 17.6

447-2L 982 (18001 207 (30) 69.9 33.2 38.9

447-6L 982 (18001 207 (30) 81.8 17.2 24.1

_-" 428-IOL 1038 (1900) 131 (19) 158.3 25.3 38.5

447-4L 1038 (1900) 131 (19) 229.3 30.3 73.5 _

447-7L 1038 (1900) 131 (191 100.2 3.5 7.3

:- 447-IL 1093 (2000) 103 (151 77.3 10.6 39.1

".: 447-3L 1093 (2000) 103 (15) 69.1 23.6 64.7

"

_# "1260"C (2300"F1 for 2 hours with _rgon Quench, plus 982°C (1800"F)

for 5 hours, pllls 871°C (1600"F1 for 20 hours.

TABLE 27. STRESS-RUPTURE TEST RESULTS ON 1.8 HM (0.070 INCH) DIAMETER

LO.NGITUDINAL TEST SPECIMENS MACHINED FROM SINGLE-CRYSTAL

: TURBINE BLADES OF HEAT-TREATED* ALLOY 4+ZE. ".

Rupture

Specimen Temperature Stress Llfe Elongation R of A

Number °C (°F) rMPa (ksl) (Hours) (%) (%) Remarks

• C

431-2L 760 (1400) 724 (1051 101.5 10.5 17.6 -

B 431,9L 760 (1400) 724 (105) 131.3 12.1 27.0

413-IL 982 (1800) 207 (30) 45.1 3.8 7.3 Casting Porosity '_

431-4L 982 (1800) 207 (301 80.4 28.8 37.8

e

"l,_ 413-2L 1038 (1900) 131 (19) 82.1 14.5 38.5

" • 413-3L 1038 (1900) 131 (19) 71.9 14.3 30.0

; [ 431-7L 1038 (1900) 131 (19) 92.7 16.5 39.4 _

L" 413-5L 1093 (2000) 103 (15) 26.5 11.8 34.1

_ 413-7L 1093 (2000) 103 (15) 24.5 15.9 41.8 _..

: *1260°C (2300"F1 for 2 hours with Argon 0uenoh, plus 982eC (1800°F1 for 5 hours, 1

,_ plus 871¢C (1600"F) for 20 hours, i!

_ _,_,

;. 83 •'
,I, !

!

,, ORIGINAL PhGZ [S

OF POOR QUALITY

TABLE 28. STRESS-RUPTURE TEST REeULTS ON 1.8 MM (0.070 INCH)

DIAMETEK LONGITUDINAL TEST SPECIMENS MACHINED FROM

SINGLE-CRYSTAL TURBINE BLADES OF HEAT-TREATED _ '

ALLOY 5*Zr.

I. Rupture

t Specimen Temperature Stress Llfe Elongatlon R of A

Nt_ber "C ("F) MPa (ksl) (Hours) (t) (%)

432-6L 760 (1400) 724 (105) 156.3 14.0 17.2

• 432-9L 760 (1400) 724 (105) 99.6 6.8 22.9 t

i.- 414-IL 982 (18001 207 _301 50.8 15.0 32.3

.!,

-_ 414-3L 982 (1800) 207 (30) 54.7 14.4 20.1

"i. }

414-4L 1038 (19001 131 (191 78.3 12.5 31.7

414-6L 1038 (1900) 131 (19) 102.2 11.9 22.9

432-5L 1038 (1900) 1_1 (19) 107.2 22.1 41.2

414-8L 1093 (2000) 103 (15) 39.0 9.5 27.8

? 414-IOL 1093 (2000) 103 (15) 32.7 10.5 24._ ,

*1260"C (2300"F) for 2 hours with Argon Quench, plus 982"C (I800"F)

_: for 5 hours, plus 871"C (1600°F) for 20 hours.

>

- Effects of slow cooling from the solutioning temperature were

reductions in tensile strength and rupture life as illustrated for

; SC Alloy 1 in Tables 29 and 30, respectively. As discussed pre-

: viously, these reductions were associaued with a significantly '

b

coarser gamma-prime size (se_ _igure 30). i

,4

, 5.3.10 SC Alloy Selection i

Based on stress-rupture and tensile data, casting yields,

_, alloy composition, and microstructural stability, SC Alloys 1 and 3 !

I

_, were selected for design property characterization and component

fab ication for engine evaluation.

SC Alloy i, subsequently renamed SC NASAIR i00, was selected

on the basis of its high stress-rupture and tensile properties,

-" plus the best yield in the casting trials. Although the mu phase

• 84
|

' ORIGINALP_G_ _

OF POOH QUALmr

," TABLE 29. TENSILE TEST RESULTS ON 1.8 MM (0.070 ZNCH) DIAMETER TEST

, _ - SPECIMENS MACHINED FROM SC ALLOY 1 TURBINE BLADES THAT

WERE IMPROPERLY HEAT TREATED.*

_ Specimen Master Temperature UTS 0.2% Y.S. Elongation R of A

:",t_. Number Heat Orien_atlon C° ("F) MPa (ksl) MPa (ksi) (%) (%)

-- 427-2L-2 VF-154 Longitudinal RT 876 (127) 827 (120) 12.8 18.3

_%• 508-2L-2 VF-154 Longitudinal RT 924 (134) 827 (120) _7.4 23.0

_: 427-7L-2 VF-154 Longitudinal 760 (14001 1131 (1641 1076 (1561 i0.i 15.3

7"

,_ 509-IL-2 VF-154 Longitudinal 760 (1400) 1220 (177) 1179 (171) 8.6 15.7

,_ 427-IL-2 VF-154 Longitudinal 871 (1600) 1034 (1501 910 (132) 14.7 18.7

508-IL-2 VF-154 Longitudinal 871 (1600) 1020 (148) 910 (132) _5.2 21.8

-- 427-5L-2 VF-154 Longitudinal 982 (1800) 648 (94) 517 (75) 25.5 31.4 .:

509-7L-2 V[.-15_ Longitudinal 982 (1800) 731 (106) 621 '90) 15.3 20.8 _/

"_ 427-3L-2 VF-154 Longitudinal 1093 (2000) 434 (63) 352 (51) 31.7 49.7

_ 509-5L-2 VF-154 Longltudinal 1093 (20001 407 (59) 338 (49} 36.0 49.4

.-. _Heat treatment consisted of 1316"C (2400°F) for 2 hours, plus 1324"C (2415°F1 for

2 hours with Argon Quench, plus 982°C (1800"F) for 5 hours, plus 871°C (1600°F) -_ ,:

i f(o19r00°20F)hourress.ulting Fan cinoolling ower wasstrength not activ thanatepdrevio untuisllyL.madeaessurhaedd. cooled to i038°C _ " """

Z"

_:" was of concern based on previous experience in other alloy systems

: with sigma, the mu level present in NASAIR 100 was a small volume _

fraction and did not have any significant effect on the properties

evaluated. :t

SC Alloy 3 was .,;elected on the basis of microstructural

stability--e.g., freedom from the mu phase and good castability.

*, The stress-rupture capability of SC Alloy 3 was also a significant "

improvement relative to that of DS Mar-M 247, but was less than SC

: NASAIR i00. ;"

85

t*

,w i
7

L

!

ORIGINALPAGEiS

OF POOR QUALI'rN'

TABLE 30. STRESS-R_PTURE RESULTS ON 1.8 MM (0.070 I_CH)

DIAMETER SPECIMENS MACHINED FROM SC ALLOY 1

- TURBINE BLADES THAT WERE IMPROPERLY HEAT

TREATED •*

i" Rupture

! Specimen Temp. Stress Llfe Elongation R of A

• Number Orientation eC ("F) MPa (ksi) (Hours) (%) (%)

..

427-7L-I Longitudinal 760 (1400) 827 (120) 40.3 8.4 38.3

!i 427-7T Transverse 760 (1400) 8_7 (120) 20.4 8.9 26.0

508-IL-I Longitudinal 871 (1600) 400 (58) 217.3 16.7 31.5 i_

!

508-IT Transverse 871 (1600) 517 (75) 40.7 12.3 28.2

_" 427-5L-I Longitudinal 982 (1800) 172 (25) 1085.0 11.9 25.6 :

509-7L-I Longitudinal 1038 (1900) 131 (19) 192.1 12.8 38.5

d

• 427-3L-I Longitudinal 1093 (2000) 103 (15) 121.2 16.2 33.0

b; 509-5L-I Longitudinal 1093 (2000) 90 (13) 202.5 12.6 33.7 _'

• *Heat treatment consisted of 1316"C (2400°F) for 2 hours, plus

1324°C (2415°F) for 2 hours with Argon Quench, plus 982"C (1800"F)

for 5 hours, plus 871°C (1500eF) for 20 hours. Fan cooling was not

_ activated until blades had cooled to I038"C (1900°F} resulting in

lower strength than previously measured.

_4

I

• 86

I

........ i
SECTION VI ._

- 6.0 ALLOY PROPERTY CHARACTERIZATION

i. Two SC alloys, NASAIR 100 (formerly Alloy i) and Alloy 3, were

characterized for mechanical, environmental, and physical proper- ?

v

F ties to validate the final design of the turbine blade. This was

. accomplished by casting additional turbine blades and test bars at

_ . Jetshapes, followed by comprehensive testing of castings and test

_. specimens by Garrett and independent laboratories.

¢

:" 6.1 Test Material Production _G

6.1.! Master Alloys _

t e

" -_ Seven hundred and twenty-five kilograms (1600 pounds) of :

? NASAIR i00 and 297 kilograms (655 pounds) of Alloy 3 were procured "_

:_ from Cannon-Muskegon to be used for design data generation and com- _ _-?

ponent manufacture (Section VIII). These alloys were provided in _ '

the form of small 120 to 150 kg (265 to 331 ib) master heats. Comparisons of master melt and alloy specifications are provided in _ .

Tables 31 and 32 for NASAIR I00 and Alloy 3, respectively. Inspec- * " i

, tion of these chemistries indicates that all of the heats ,_ere

i

acceptable. In addition, a portion of NASAIR i00 (Alloy !) heat _ ;

.- VF-154 remained from the alloy developmen task and was used to . ;

, cast three molds; consequpntly, this heat chemistry is also .:. t

' included in Table 31. _

6.1.2 Blade and Specimen Castings _ i

• t, _

_'< SC TFE731 blades, Ears, and rectangular slab castings were _, !

_ produced--using the casting technology described in Section III--to ,_ t

provide test material for the SC alloy property evaluations. I

_ 87
ORIGINALPAGE t3

OF POOR QUALITY

TABLE 31. NASAIR i00 MASTER MELT CHEMISTRIES (WEIGHT PERCENT).

Heat

. Mo 0.8-1.2 0.0 1.0 1.0 1.0 0.9 1.0 1.0 1.0

W i0.0-II.0 10.4 i0.2 10.3 10.4 10.5 10.4 10.4 i0.5

Ta 3.1-3.5 3.4 3.3 3.4 3.3 3.3 3.3 3.3 3.3

Ti 1.0-1.4 1.1 1.0 1.0 1.0 I.I 1.1 1.1 1.14

C_ 8.5-9.5 9.0 9.1 8.8 8.0 0.7 8.0 8.9 9.0

Co 0.5 max -- <0.1 <0.1 <0.1 <0.1 <0.15 <0.1 <0._

A1 5.6-5.9 5.8 5.7 5.6 5.7 5.6 5.8 5.7 5.7

Zz 0.03 max 0.01 <0.01 <0.01 <0.01 <0.01 <0.01 <0.01 <0.006

B 0.002 max -- <0.002 <0.002 - - 0.0018

C 0.006 max 0.0034 <0.0025 <0.003 <0.003 <0.005 <0.006 - <0.0002

Ni Balance Balance Balance Balance Balance Balance Balance Balance Balance

. WOTE: Cannon-Muikegon Certification Chemical Analysis

L ._:

2-

;_. _

TABLE 32. ALLOY 3 MASTER MELT CHEMISTRIES (WEIGHT PERCENT).

" I Alloy 3 Heat

; Element I Specification VF-291 VF-292 VF-442 ;i

Mo 0.65-0.95 0.8 0.8 b.8

W 9.0-i0.0 9.5 9.5 9.5 !

Ta 3.1-3.5 3.3 3.3 3.3

" Ti 0.9-1.2 1.0 1.0 1.0

C= 8.0-9.0 8.6 8.6 8.6 ;

Co 4.5-5.5 5.0 5.0 5.0

A1 5.5-5.8 5.8 5.8 5.6 Ii

Hf 0.2-0.3 0.25 0.25 0.24

!

C 0.006 Max <0.003 <0.003 0.0026

t i

, B 0.G02 Max HA NA 0.002 I

i

Zr 0.03 M_x <0.01 <0.01 <0.01 t

I Ni Balance Balance Balance Balance

NA : Not analyzed. 1

!

i

. ;

f

' 88
Results from this casting 'campaign are reviewed in Tables 33

through 35 for SC NASAIR 100 and SC Alloy 3 TFE731 blades, bars,

: and burner-rig specimens, respectively. These tables indicate that

_ of the blades, bars, and burner-rig specimens that were filled with

i" metal, yields of material suitable for testing were 49, 63, and

::i- 61 percent, respectively.

" Compositional analyses of SC NASAIR 100 and SC Alloy 3 blades

I : cast during the campaign are provided in Tables 36 and 37. All of

!. the blades analyzed were either within or very close to the specification limits. The minor deviations indicated are not considered

' significant. ,_

TABLE 33. VISUAL INSPECTION RESULTS FOR SC TFE731

_'_ B_.DES OF NASAIR 100 AND ALLOY 3.

Blades Filled Single-Crystal ":""

_= " Alloy Mold No. with Metal Airfoil and Root* '

._° ,. ........... ) l

NASAIR i00 427 i0 o

457 2 1

, 458 12 8

• 508 3 1 i

: 509 i0 3

511 12 3

512 i0 4 .

_. 513 I0 4

591 12 7

647 12 2

648 12 9

|

. ._ Alloy 3 675 I II

_ % 688 12 9

_: 4 691 I I0 2

_[O01J grain orientation within 15" of the blade

axis. {%

, 4

89 _
OF pOOR Q_L__

TABLE 34. VISUAL INSPECTION RESULTS FOR SC BAR CASTINGS.

SC

_ Bars Pilled Bars Acceptable

. Alloy Mold No. with Metal for Fatigue Testlnq _

' NASAIR I00 521 I_ 2

544 12 12

547 6 3

552 12 8

553 12 7

_67 12 I0

, 568 12 10

._ 569 12 11

_. 570 12 10

653 12 4

A11o¥ 3 582 12 7

666 12 5 '_

687 7 5

:' " 689 7 7

_. 695 12 6

7

709 12 10

i

*[001] grain orientation within 15" of the bat axis.

; TABLE 35. VISUAL EX_J4INATION OF BURNER-RIG SPECIMEN CASTINGS.

_4

I Specimens Filled Single-Crystal Alloy Mold No. wlth Metal Specimens

NASAIR 100 548 i0 S !

NASAIR 100 549 12 8

|'.

" t

_ , Alloy 3 655 12

AIIoy 3 673 12 8

I

9O
!

!

, ORIGIN/_t._;:,

OF POOF: ("_,:A,-_i_'',

During the casting campaign, test material was selected bl,

etching the blades and bars and visually examining them for [00i]

orientation within 15 degrees of the bar or blade axis. A lim/ ted

number of confirmatory tests for orientation were conducted using

!

the Laue method at Technology of Materials in Santa Barbara,

2

California. Laue inspection results, provided in Table 38, con-

,:? firmed the orientations determined by visual exan.ination in all

o"

i cases.

.[.

TABLE 36. CHEMISTRIES OF SC NASAIR 100 TFE731 TURBINE BLADES.

"o COMPOSITTON fWEIGRT PERCENT)

• , 427 427-8 VF-154 • )6 _ 0.I 5.6 <0.01 <0.002 i 0.004 Bal

"_ 508 508-3 VF-154 )6 _9.02 0, _ _ 8 <0.01 <0.002 ]0.003 Bal '_

_" , 509 509-8 VF-154 )4 I 9.24 " 3.7 <0.01 <0.002 ) 0.006 Bal

:" :' 511 5n-1 _-220 (0.99 10.5 3.2_ )0 _6 J9.87 .1 3.5 <o.ol <0.002(0.005 Hal "

, 513 513-10 VF-220 J 0.97 10.4 3.30 0. )8 9.97 .I ).6 <0.01 <0.002 0.005 Bal

_: ', 593 591-30 VF-235 I0.80 10.2 3.36 1. ,z 9.68 .i 3.7 <0.03 <0.002 0.004 Bal

"'" 647 647-10 VF-235 ( 0.94 10.6 3.50 I. }5 9.21 ,I _.6 <0.01 <0.002 0.00_ Bal

• : 648 648-11 VF-2_5 I 0.82 10.5 3.52 I. )0 9.50 .1 5.6 <0.01 0.002 i 0.003 Bal

• _,.o_,,- ._o. I 0 _ lO.O 3.1 1.0 I 8.510.5I .6 0.03 o.0o2Io.o06I

:. Ranoe Max. _ Max.

Analys_s performed by the Garrett Turb,ne Engine Company. ."

: i

TABLE 37. CHEMISTRIES OF SC ALLOY 3 TURBINE BLADES.

B

,.OMPOSITION (WEIGFT PERCENT)

Mold No. Blade No. Heat No. Mo W # Ta I Ti Cr Co A) Hf _r B C I Ni

. . I I .

_Y ( 675 675-11 VF-292 0.78 ,s 0.96 I8.20 5. 5.5 0.27 l_ <0.01 II1<o.oo2,o.oo6 IPal

I I I I IBal

688 688-10 VF-292 0.76 I 9.4 3.20 0.90 8.90 %. 5.6 0.30 i <0.01 1<0.002 0.006 ,

---, i ¢ [ , I I ' "

°°°1°°I°I: Sp,-':i- Ntn 0.65 9.0 I 3.1 0.9 I 8.0 4. 5.5 0.2 0.03 I _ i

f _cation B 1 ,

Bang_ Max 0.95 lO.O ) :).5 1.2 I 9.0 5. 5.8 0.3 ) Max Nmx Max

Analysis _erformad by The Garrltt Turbine Engine C;mpany.'

' 91

1®t
i'

i

r_

; r

OE POOR QU A_'T_I

Visual O_ientation Laud Orientation

Alloy of Blade* .... Blade* MFB Specimen**

b

SC NASAIR i00 -<15 G of [001] 12" of [001]

SC NASAIR i00 _-15" of [0011 14" of [001]

J'_ SC NASAIR i00 -<15 e of [001] 3.2" of [001]

i SC NASAIR i00 >15" of [001] 31" of [001]

', SC NASAIR i00 >15" of [001] 134" of [001]

SC NASAIR i00 -_15" of [001] 1.5" of [001]

'_- sc HASAZRi00 -<1°5 of [0011 8° of [0011 i

SC Allog 3 -<15° of 1001] 15.5" of [001]

d

SC Alloy 3 -_15" of [001] 4" of [001] ,=

- SC Alloy 3 -_15" of [001] 8 ° of [001] "_

_ *Orientation from the blade stacking axis.

_ **O_lentation from the specimen axis.

?: i .....

t To asses_ the potential significance of mu phase precipitation, fully heat-treated SC NASAIR 100 and SC Alloy 3 blades and

; bars were exposed at 982°C (1800°F) for 1000 hours.

. Metallographic and scanning electron microscope (SEM) exami-

_ nation (Figure 40) to asses:_ the effect of exposure on NASAIR i00

oJ

and Alloy 2 microstructural stabilit2 indicated that the major

effect of 1000-hours exposure on both alloys was garma-prime coarsening, with NASAIR 10O also exhibiting some acicular pha;_e (mu) in ,-

the dendritic areas. Alloy 3 exhibited no evidence of topologi- i

;_ tally close-packed phases, i

Extraction and identification of secondary phases was perfor- ;

med by Dr. J. F. Radavich of Mic:omet Laboratories in West

Lafayette, Indiana. Results of this evaluation, provided in

Table 39, and indicate that the followinq secondary phases were

" observed in addit. _n to _.he gamma and ganu,,--_rime majority phases:

92
' P72137 m -9

• _ Mu PRECIPITATION 7' COARSENING ;

' 7 COARSENING ,

5000X MAG

FULLY PROCESSED+ 982 C (1800 F)/1000 HOURS

", Figure 40. Furnace Exposure Effects on SC Alloy :

,!

- Microstructure. ,,

':' 93

I
®

TABLE 39. RESULTS OF SFCONDARY PHASE EXTRACTION AND

CHARACTERIZATION OF SC NASAIR 100 AND SC

ALLOY 3 TURBINE BLADE CASTT_GS.

, Relative Amount of Phases

:i ,| Allo_ Cond i t ion MC T iN _W mu

SC NASAIR !00 HT* Weak Weak Strong --

J_

:, SC NASAIR i00 HT _ Weak Weak Strong Strong

l: +982Oc (1800OF)/

! i000 hrs.

t

SC Alloy 3 HT** Weak Weak .... ;

SC Alloy 3 HT** Weak Weak Weak --

[ +982°C (1800°_)/ _

1000 hrs. _

b" L

HT* - Neat treatment consisted of 1Z_6°C (2400°F) for 2 hours,

.w plus 1324°C (2415°F) for 2 hours with rapid Argon Quench,

plus 982°C (1800°F) for 5 hours, plus 871°C (1600°F) for 20

hour s.

i

HT** - Heat treatment consisted of 1288°C (2350°F) for 2 hours,

pJus 12£6°C (2365°F_ for _ hours with rapid Argon Qdench,

plus 982°F (180D°F) for 5 hours, pluF @71°C (1600°F) for 20

_ hours.

o SC NASAIR I00 - _W was the p_'incipal secondary phase that

was detected in the fully processed micrestructure; t_'ace

amounts of MC carbide and TiN phases were also d _._ected.

AfteL the 1000-hour exposure at 982°C (1800°F), _be n,u

phase was also identified as a principal secondary phase.

o ,_C A.l.loy 3- In t_u fully processed condition: only trace

amounts of the MC carbide and TiN phases were detected.

Altar the 100C-hour exposuL'e at 982°C (1800°F), trace

amounts of ,_W wer_ also identified. No mu phase was

observed.

\

94
Following casting and inspections, blades and bars were heat

treated, machined into test specimens, and tested. Results are

pr._vided in the following paragraphs.

6.2 Mechanical Properties

SC K.S_IR i00 was characterized for tensile, creep-rupture,

n!%h-cyclr fatigue (HCF), and low-cycle-fatigue (LCF) strengths.

S_m:lar data were obtained for SC Alloy 3, but in reduced quantities. These evaluations are reviewed in this section.

6.2.1 Tensile Properties

SC NASAIR 100 _

' Tensile data for SC NASAIR 100 were obtained by Joliet Metallurgical Laboratories (Joliet, Illinois) over the range of 21 ° to

_ I093°C (70 to 2000°F), using 1.8 mm (0.070 inch) diameter minibars

that were machined from TFE731 turbine blades (Figure 13). Blades

tested in this evaluation were produced from four master heats.

Test data for uncoated specimens are provided in Table 40 and

average properties are shown graphically in Figure 41. Inspection

of these data indicates that strength is maximum and ductility is

minimum at 760°C (1400°F). The test results also indicated comparable tensile strengths for longitudinal and transverse specimens, provided orientations are within 15 degrees of [001] and

within 15 degrees of [i00], respectively. For transverse specimens

_: with axes exceeding 15 degrees of the [i00], strengths were gener-

!i ally lower.

Variability of tensile strength was also maximum at 760°C

(1400°F) as shown in Figure 42. The amount of scatter observed is 1

very similar to that found in SC Mar-M 247 (Figure 14). Based on I

C - 9s

}
ORIGINALPAGE I$

OF POORQUALITY.

TABLE 40. TENSILE TEST RESULTS FOR 1.8 _N (0.070 INCH) DIAMETER TEST SPECIMENS MACHINEI!

FROM SC NASAIR i00 TURBINE BI _DES.

Ortsntot_on

M/thLn ] WLth_n

,,o,.. ,Nulbo[ ,,..*.r Beat Orlen_rat .... ;on 1o__0(001} .,,... 1I 150, o_ |1001..... 15Of +.,,... (100l .:_.,_;;a 141Du. I (kli) NPI 02, ,9(kSi) .oo_.*,.n (I) . o, (I)

511-8L-2 VF-220 r,_ngitudLnal 9 21 (70) 972 61413 945 (1373 13.0 17.9

511-12L-2 VlP-220 Long_tudAnal Z 21 (?01 1027 61493 972 61413 13.2 19.0

64_-IL-2 VF-235 Longitudinal X 21 (70) 1027 6148) 986 (1433 16.0 21.9

648-3L-2 Vf-235 LoflgLtudLnol x 21 (70) 1186 61723 1034 (150) 9.8 11.8

591-11L-2 VlP-233 Longitudinal X 21 (70) 1007 (1463 _86 (1433 16.1 22.1

148-37-2 Vr-235 Transverse X 21 6703 986 61433 IS0 (1393 19.5 28.2

511-77-2 vr-220 TransverSe X 21 (70) 956 6139) 917 61333 15.3 1_,1

513-67-2 VF-22d Transverse X 21 6703 903 (131) 869 61263 19.1 22,4

64&-2L-2 Vt-235 Longitudinal X 649 61200) 1138 6165| 965 (1403 8.6 9.7

648-4L-2 VF-23SJLonqLtudLnal Z 649 612003 11110 61613 958 (139) 11.6 12.5

591-15L-2 'F-235 LongLtud/_sl X 649 "1200) 1234 (1793 1034 6150| 5.8 6.9

648-47-2 VF-235 Transverse X 649 61200) 1078 61563 965 (1403 5.4 8.3

648-67-2 VF-235 Ttanevarna X 648 (1200) 1089 (1583 938 6136) 10.0 11.8

668-77-2 VF-235 Trannve[Is X 649 (1200) 1062 (154) 945 (137) 0.? 19.8

401-3L YP-154 Longt¢adtnnl X 760 (14003 11379 6200) 1289 6187) 3.7 11.8

481-7L VP-IS4 [Longitud/nsl X 760 (t400) 11407 (204) 1338 (1943 6.0 14.2 '

481-1L-1 VF-154 I LongLtudLnsl X 760 614003 1358 (1973 1276 (185) 3.6 9.9

513-8L-1 vr-_u !LOngLtu_qnl X 760 (1400) 1303 (109) 1096 (159) 7.3 9.0

.'. 513-8_-2 v?-220 _onLtudinal X 760 (1400) 1262 _1833 1089 (158) 2.9 S.5

648-5L-2 Vr-235 _ongitudinal _ X 760 61400) 1255 6182) 11124 61633 2.5 2.7

648-6L-2 VF-235 LongLtudinal [ X 760 (1400) 1420 (206) 1338 6194) 4.4 5.4

591-9L-2 Vr-235 LongLtudtnal I X 760 (14003 1358 6147) 1303 61893 7.1 10.4

753-1L-1 Vr-237 _ongitudinal X 760 61400) 1379 (200) 11269 (104) 6.0 12.7

. .;

765-1L-,1 Vf-237 LongLtud/nal X 760 614003 13"9 (200) 1234 (179) 6.7 14.9

7?4-1L-1 _F-237 LongL_udLnal X J760 614003 1351 (196) 1220 (1773 1.7 4.1

513-87-1 VF-220 Transverse X 760 (1400) 1262 61833 1117 61623 5.7 7.8

513-8T-2 Vr-220 Transverse X !760 (1400) 1255 6182) 1136 6185) 6.0 10.4

511-12T*2 v]r-220 Transverse X 760 (1400_ 1110 (161) 098 (110) 11.0 13.8

481-1L-2 vr-154 Longltud|n_l X 871 (16003 1124 (161) 925 6134) 7.5 15.3

457-1L°2 V_-154 _ong/tudinsJ X !871 (16003 1124 (1633 988 (143) 13.3 17.4

511-11L-2 W-220 Longitudinal X 871 616003 1117 (1623 965 _'¢_) 11.2 14.6

_]-?L-2 _F-220 _ons(_ud|nal X 871 (1600) 1151 (167) 1000 (14_) 11.3 12.0

551-19L-1 Yr-235 LongL_udfnal X 871 613003 972 61413 $83 [1283 11.9 10.7

5_1-19L-2 W-235 LongLCudLnal X 871 (16_:) Jd6 614)) 903 6131) 8.5 17.4

511-117-2 V?-220 Transverse X 871 616003 952 613_) 821 (119) 15.4 17.0

511-7T*2 VP-220 T_snfva[ae X 871 (:600) 1083 6157) 1000 (145) 4.9 9.9

513-97-2 vr-220 Tranavar|s X 871 616003 1014 61473 9]9 (136) 13.7 17.9

481-2L-2 V_-154 _ongitudinal X 982 (3800) 862 6125) 648 (94) 17.4 18,3

_11-10L02 V_-220 Longitudinal X 982 (18003 821 61193 717 6104) 17.9 25.1

511-9L-2 VF-220 Longitudinal X 882 (18003 827 (1203 738 61073 17.3 25.3

648-7L-2 Vf-235 LongLtudLnal X 862 (18003 841 (1223 731 (1063 23.5 36.5

591-2LL-2 _-235 _ong_ud_nsl X 982 618003 827 6120) 696 6101) 26.8 28.6

551-22L-2 W-235 Lon0Ltudinal X 982 (ll00) 862 61253 724 (105) 22.5 31.9

• 4|1-2L-1 V_-154 LongLtud_nsl X .08] (2000) 483 (70) 386 (56) 20.5 39.3

458-9L-) Vf-154 LongL_udtnel X .083 (2000) 510 (74) 400 (58) 34.1 50.1

_** 511-7L-2 VF-220 Long/_udLnnl X .OSJ (2000) 469 (68) 37_ (55) 36.1 48.3

513-6L,2 Vr-220 Longl_ud_nal X .08) (2000) 469 (68) 365 653) 26.8 38.0

648-8L-2 vr-235 Longs_udin81 X .093 (2000( 462 (673 386 (56) 26.5 52.2

581-23L-2j V'lF-235Long_udtn81 X ,091 (2000) 483 (70) 338 649( 29.4 55.0

511-9T-2 Vr-220 Trsnsvaraa x 093 (2000) 448 (65) 345 (503 21.1 25.3

511-L07-2 VW-220 Trnnaverao X _09] (20083 441 (64) 328 ,49) 23.8 29.4

5_1-97-2 V_-235 Transvsrnn X 083 (2000) 455 (66) 310 (45) 27.0 49.4

.... j

lOLl TCeltJleOt aO_llltOd of 1316°C 12400+P) fOr 2 houri,

pl,l 132_°C (24150P1 for 2 hour| wXth Rapid Acgon Ouen+h_

plus 912 C (18000f) for 5 houce plus 871"C (1100"_1 (oc

_0 bouts.

96
/

ORIGINAL PP,G;£ I$

OF POOR QUALITY

1000(210)-

1000 (145)- __

= \

T

' _ 500 1731

40 _ "_

_ ELONGATION 20 _ _ "'

• !

0 500 1_0 °C ":

i • w i !

0 500 1000 1500 2000 OF

TEMPERATURE

"2

Figure 41. Average Tensile Properties of SC NASAIR i00 with _

Specimen Axis within 15 ° of <O01>. ._

_ a

the results for SC MAR-M 247, about half of the 250 MPa (36 ksi) ?

width of the =catterband is attributable to permitting specimen ._

; orientation to vary within 15 degrees of the <001>.

_L

Tensile tests were also performed on diffusion aluminide _

._ (RT-2I) coated 1.8 mm (0.070 inch) diameter minibar specimens. The

; test specimens were machined from solution heat-treated SC _

NASAIR i00 turbine blade castings, and then RT-21 coated by Chrom- .;'

alloy American Corporation, Orangeburg, New York, to a coating ._

thickness of 43_m (1.7 mils) followed by an 871°C (1600°F), 20 hours [_

post-coat heat treatment. Tensile test results from room tempera- '..

ture to 1093°C (2000°F) for longitudinal and transverse orienta- i!

tions are provided in Table 41. The effects of coating on tensile

properties were reduced ductility up to 760°C (14C0°F) and

97
ORIGINALPAGF-iS

OF POOR QUALi'i-Y

1370(200)_

_::.:.V.."..'_::::::::_.::::.._k •.':"_2_._

:....:_,_>'_..,'..:: _:..;....

,_ ,,.,,,.... ....:.._,...:,.._..,.::".."._'...-'% _,,; ,,, ,,,:,,:, :,,,:,, ,,/,,,,,,,,,'.__,j_. -v,-. ,-:-,

•"" ' "":".:+.";_.:"""!':_"." ":'_:" " '"" '"'" "_k

1034(160).... : "" :" ""..... :".. ": _.

,,._,,_..:.:'-.....p.... ::. •

..... .... ..:

i" Li!"___& "''_"" " "_ " "_'_"":_':::':"

uJ

"%

leo (1001-

3451501.

_ SPECIMENAXIS WITHIN 15°F<0011 _.

)4 SPECIMENAXIS >15°F <001>

\o

0 ! • • i i ! ! T ! "!

536(1000) 816(1500) 1093(2000)

.- TEMPERATURE,°C(OF) -" '--

Figure 42. Yield Stress Data for SC NASAIR i00.

TABLE 41. TENSILE TEST RESULTS FOR DT_FFUSION-ALUMINIDE (RT-21) COATED

1.8 MM (0.070 INCH) DIAMETE_ TEST SPECIMENS MACHINED

FROM SC NASAIR i00 TURBINE BLADES.

Of _ent_g_on

NLghzn I_Lth|n

Spec_Lalen i Nagger Graxn 15 Settees 15 Deggeee >15 Degree| Temperat:ure UT$ 0.24 ¥.S. glonqetLon R _ A

,.NUlllber Heat Orlentet:Lon od_ (OO|J og (100( o( (1001 *C _*r) J4JPi (kl; MPe (ksx) (_l) (t)

511-01,-2 vP-220 I_9_tud_nal x 21 (701 1172 (170) 1000 (1451 6.8 6.2

513-fi[,-2 VIP-220 t.ongl BudLnal X 21 (701 1117 (162) 1000 (1451 9.4 4..!

" 1911-4T-1 Vr-220 Trensve_aie x 21 (70) 89_ (130} 083 (1281 9.1 20.9

1513-1T-1 VF-220 Transverse X 21 (701 810 (132 889 (1381 1.1.8 17.8

1511-61-2 VIP-220 [4)or | t t;_lAnal ]( 649 (1_00) 1165 (168) 1007 (146) 4.4 ill.'_

513-4[,-2 V1_-220 LOOg_LgudLnm2 X fi48 (1200) 1124 (163: 10i4 (147) 5.3 7.5

511-5T-1 VP-220 Trensveree x 648 (12001 1138 (155) 1014 (1471 2.5 2.8

' 513-3?-1 vr-230 Transverse X _48 (1200) 1055 (1831 858 (1._8} 2.2 4,8

t,

511-81-3 VP-2_0 L_O_gLt u(I _nal X 648 (1400) 1330 (1841 1220 (177) 0.S 0.7

_-_ :513-31-2 VP-230 LcmgAtudLnal X 649 (14001 1310 (1901 1279 (1711 3.0 4.7

._ d 511-0T-1 VIP-220 Transverse X 449 (1400) 1282 (106) 1240 (101} 2.2 _,3

.! S13-5T-1 VP-230 Trensvcree X 048 (1400) '._Sq (1741 1124 (143) 3.5 4.7

Sl1-41-2 VF-220 5o_g[tud/nal X 871 (1400) 1103 (1_0) 878 (1421 1i.7 18.5

S13-2S,-2 VF-220 f,o_gLtudinsl X 011 (16001 1117 (1521 879 (1421 15.1 15.?

511-1&-2 VIP-220 LongL gudLnal X 108) (2000) 441 (_4) )31 (48) 31.1 48,6

'913-11-2 Vr-220 Long / tudJnal x 108:) (2000) 414 (40( 303 (44) )1.8 44.2

$11-3T-1 VIP-320 Transverse X 108] (2000) 427 (_2) 310 (45) 23.8 34.8

81]-2T-1 Vr-220 Transverse X 10_13 (2000) 434 (63) 331 (40) 4,7 14.8

Beat treaWiang (:onmtsted o_ 1314"C (2400.r) _or 2 hours.

,L plus IJ24': (2419*r) _or 2 hours vith Bap_d Ag4on Ouenc:h,

7 I?lUe BT*;II soared, plus 171 C (1500 IP) I_or 30 hours, , ,

98
ORIGINAL PAGE 15

OF POOR QUALITY

,, increased yield strength at room temperature. These differences

are not considered significant since the blade root will not be

coated.

Tensile test results for SC NASAIR 100 in the fully processed

+ 982°C (1800°F)/1000-hours exposure condition are provided in

Table 42. Inspection of this data indicates that, in general, the

• exposure reduced the tensile strength of NASAIR 100 while slightly

: increasing ductility. At 871 ° (1600 °1 and 982°C (1800°F), average

yield and ultimate strengths were reduced by about 15 percent. ._

At room temperature, the 982°C (1800°F)/1000 hours exposure

0° resulted in strength reductions in the order of 30 percent, while

• increasing ductility by about 60 percent This lo&s in strength,

., however, was not considered significant in the turbine blade since

-s -

TABLE 42. SC NASAIR i00 TENSILE TEST RESULTS AFTER EXPOSURE AT

982°C (1800°F) FOR i000 HOURS.* 4

I I :

Specimen IMaster Temperature UTS 0.2% ¥.S. Elongation 2 of A

,, Number I Heat Orientation** C° ("F) MPa (kei) MPa (ksi) (%) .. (%1

647-9L-2 VF-233 Longitudinal 21 (701 750 (II0) 676 (981 21.0 28.1 _,|

647-3L-2 VF-233 Longitudinal 21 1701 765 (Iii) 600 (87) 26.7 28.1

591-36T-I VF-235 Transverse 21 (70) 848 1123) 662 (96) 30.6 31.2

.j,

591-36T-2 VF-235 Transverse 21 (70) 834 1121) 655 (95) 31.5 31.2 r

R

',* 591-36L-2 VF-235 Longitudinal 871 (16001 952 1138) 793 (1151 12.9 17.2

'_ 648-95-2 VF-235 Longitudinal 871 (1600) I000 (145) 855 (124) 27 0 26.5

647-125-i VF-233 Longitudinal 982 (18001 703 (I02) 614 (89) 25.0 26.1

591-3sL-2 Longitudlnel 9,2(1,00, ll05, 6. 1921 18. 27.8 ii

*Tensile tests were conducted using 1.8 mm (0.070 inch) diameter test specimens machined "

from SC turbine blades that have been heat treated end then exposed to 982°C (lS00*F) for ,_,

1000 hours. Prior to the 982"C (1800'F) exposure, hilt treatment consisted o_ 13_6°C "_

(2400°F) for 2 hours, plus 1324tC (2415'F) for 2 house with rapid Argon Quench, plus 982°C ._

._ **L(o1800 ngit"udF)inalfor and 5 hourtrsa,nsveplurses 871°C orlentet(lo160ns0°F) areforwit20hinhou1r5s". of [001] end 15" of [i00], respec- _

tively. :I

,"3_I'

"" )

_ 99

%
ORIGINAL PAGE |_

OF POOR QUALITY

the cooler portions of the turbine blade (i.e., firtree attachment)

will never reach temperatures approaching 982°C (1800°F).

SC Alloy 3

Comparative tensile data for SC Alloy 3 MFB minibar specimens

were obtained for the same conditions evaluated for SC NASAIR 100.

Tensile data for uncoated SC Alloy 3 specimens are provided in

Table 43, and average tensile properties for SC Alloy 3 and SC

NASAIR 100 are compared in Table 44. Inspection of these data

indicates that SC Alloy 3 is weaker in the 760" to 982°C (1400 ° to

1800°F) temperature range relative to SC NASAIR 100. Also, SC

Alloy 3 does not exhibit the pronounced maximum strength at 760°C

(1400°F) that was observed with SC NASAIR i00. Scatter in the

> strength data appeared to be comparable to that previously observed

: for SC MAR-M 247 and SC NASAIR 100. Ductility values for SC

7

_ Alloy 3 were slightly increased at intermediate temperatures relative to SC NASAIR 100.

TABLE 43. TENSILE TEST RESULTS ON 1.8 MM (0.070 INCH) DIAMETER TEST SPECIMENS MACHINED

FROM SC ALLOY 3 TURBINE BLADES.

Or lentation

Specimen 4alteL- Green L_ DegEe¢l 15 Oegreal 15 i Temperature UTB 0.2_ Y.S. 11o,oatic_l R o| A

NuBbet heat aria.titian o' (O01[ of 4100) of *C (oF) NPa (kI_) NPO (kli) (t) 44)

601-11-2 V_-Lgl Longitud_nal X 21 (70) 1117 (162) 1000 4145) 716" 16.6

615-tr-2 VP-291 I,ongi tudinal X 21 (70) %38 4136) 238 41)6) 21.3 24.1

483-7b VF-99 Longitudinal ][ 760 (1400) 1200 (174) %31 (136) g.S 14.4

Y 4$3-121 Vr-|$ nongt tud_nal X 760 41400) Ll?g 4171) 158 (119) 7.7 15.2

in-2(,-;+ Vlr-;191 Long ;¢udinal X 760 (1400) 931 1116) li] 4128) 14.4 1$,6

675-3L-2 vr-LSL Long i _ud Anal X ?60 11400) 1220 1177) 1000 4145) 2.6 S.3

$88-2T-2 VW-Lgl TranJvecae X 760 41400) i 1220 (177) 1000 4145) 3.5 5.3

675-1T-2 W-291 Trinevecne x ?O0 41400) 1014 (147) 846 4130) 7.4 12.5

7|7-11-1 Vr-292 L_noitud Lit1 X 760 41400) 1055 4153) |8$ 4129) 5.3 |.$

• 7_1-11-1 VF-292 Longitudinal ; 750 (1400) 1365 (Ltl) L00t (ISl) 6.1 14.0

{ _, 7$]-LL-L VF-442 Longitudinal X 760 41400) 111l 4194) 1234 (l?S) 4,5 S.S

• ( 680-3b-2 VF-2_I t_ngitudLnll x 871 (1600) 1111 4152) 862 4125) 10.0 17.2

675-41-2 Vr-LgL L_ngLtud,naL X 871 41600) 1103 4160) 807 4117) 14.7 24.5

618-41-2 vr*291 Longi t:ud zn41 x 982 41800) 111 4112) $21 490) 3i.1 41.3

675-51,-2 vr-291 bongi _udinal X 983 41800) J4$ 4108) 621 491) 30.$ S0.0

675-11-2 Vlf-LSl Longitudinal X L093 (2000) 414 460) 281 441) ]0,7 S0.0

$75-61-1 vr-2)L bongitudLnil X 1093 (2000) 427 ($3) ]10 ($b) ]],S S0.3 ,

675-5T-2 Vf-lJl Transverse X 1093 (2000) 421 (il) 269 (LS) 31,3 53.0

$?S-IT-2 VT-LJl f/inlverie X 11053 (LO0O) 441 ($4) ]S_ (Sl) 15,2 40.6

HOOt ¢¢eltlent oonlilted Ot _2880C (2350*t) for 2 houri, m_lllad Ln _lUqO I_[k.

_" pLUl 1296*C (2365*') tar 4 houri with Rapid Argon Ouehch.

: pLul 9020C (LO0Oer) for $ hOUri, pLLm 8710C (lt00"r) tic

20 houri.

100
ORIGINALP,"-,(_.:;3.

OF POOR QL,A:.._ITY

TABLE 44. AVERAGE TENSILE PROPERTIES OF <_0.]_ t15 ° ORIENTATION SC ALLOYS.

Orientation SC Alloys

Ultimate 0.2% Yield

Temperature Strength Strength Elongation, R of A,

Alloy °C °F MPa (ksi) MPa (ksi) % %

NASAIR 100 21 (70) 1034 (150) 979 (142) 14.6 20.5

649 (1200) 1117 (162) 965 (140) 8.4 11.5

760 (1400) 1338 (1942 ]_?.0 (177) 4.9 9.1

871 (1600) 1117 (162) 972 (141) 9.6 13.8

982 (1800) 841 (122) 710 (103) 20.9 27.6

1093 (2000) 476 (69) 365 (53) 28.6 48.3

ALLOY 3 21 (70) 1027 (149) 965 (140) 14.5 20.4 ,_

760 (1400) 1172 (170) 986 (143) 6.8 11.4 ._

, 871 (1600) 1110 (161) 834 (121) 12.7 20.8 ;

_" 982 (1800) 758 (110) 627 (91) 34.2 49.1

: 1093 (2000) 427 (62) 372 (54) 28.8 48.7

i

A limited amount of tensile tests were performed on diffusio-.

aluminide (RT-21) coated minibar specimens, and test data is provided in Table 45. As was the case for coated SC NASAIR i00, duc- "i

tility was reduced for coated SC Alloy 3 at 21 ° and 760°C (70 ° and i

1400°F). Coated SC Alloy 3 was slightly stronger at 760°C (1400°F) t

and weaker at I093°C (2000°F) than the uncoated specimens.

Tensile test results for SC Alloy 3 specimens in the fully

processed plus 982°C (1800°F)/1000-hour exposure condition are pro-

:t

vided in Table 46. Inspection of these data indicates that the t

exposure reduced strength and increased ductility at room temper- !

ature. At 871 ° and 982°C (1600" and 1800°F), ultimate strengths

were reduced, but yield strengths remained about constant or ;

increased slightly.

t

: i

]

i01
ORIGINAL PAGE i3

OF POOh QUALITY

TABLE 45. TENSILE TEST RESULTS ON DIFFUSION-ALUMINIDI: (RT-21) COATED 1.8 MM

(0.070 INCH) DIAMETER TEST SPECIMENS MACHINED FROM SC ALLOY 3 BLADES.

Speclmenl_aster I TempeEature_S0,2_J_,E1ongatlonl_of_

Number Heat Orientation C" (°F) MPa (ksi) gPa (_si) (t) (t)

62fi-7L-I VF-291 Longitudinal 21 (70) 1041 (151) 952 (138) 6.3 9.0

688-8L-1 VF-291 Longitudinal 21 (70) 972 (141) 938 (136) 12.5 15.9

675-2L-2 V¥-291 Longitudinal 760 (1400) 1158 (168) 945 (137) 3.4 5.8

688-5L-2 _-291 Longit,ldinal 760 (1400) 1220 (177) 952 (138) 4.2 8.8

675-7L-2 VF-291 Longltudinal 1093 (2000) 393 (57) 228 (33) 29.4 58.9

688-6L-2 VF-291 Longltudlnal 1093 (2000) 379 (55) 241 (35) 35.5 55.3

Heat treatment consisted of 1288"C (2350"F) for 2 hours, plus 1296°C (2365"F) for

4 hours wlth rapid Argon Quench, plus RT-21 coating, plus 871"C (1600"F) for

20 hours.

*Specimen orientations are all within 15 degrees of [001] direction.

6.2.2 Creep-Rupture

SC NASAIR i00

Stress-rupture and creep data were obtained for longitudinal

and transverse SC NASAIR 100 MFB minibar specimens (Figure 13) with

stress axes within 15 ° of the [001] or [i00] directions. Uncoated

creep-rupture data obtained in the program are presented in

Table 47. Analysis of the data indicated that there was no significant difference for longitudinal and transverse specimens as

long as the specimen orientation was within 15 ° of an _0_ direction.

102
/

ORIGINAL PAGE [_

OF POOR QUALITY

TABLE 46. SC ALLOY 3 TENSILE TEST RESULTS AFTER EXPOSURE AT

982°C (1800°F) FOR 1000 HOURS.*

Specimen TMaster I Temperature _S 0.2% Y.S. Elongation IR _ A

Number _Orientation** C ° ("F) _a (ksi) MPa (ksi) (%)

688-7L1,-2F-291Longitud n,211 (70) 731(106)724(10S) 55.3 I _

675-8T-2 t_-291 Transverse 21 (70) 917 (133 841 (122_ 16.6 20.9

: 688_9L-2 I_-291 Longitudinal 871 (1600) 979 (142 883 (128) 13.5 15.9

688-10L-21VF-291 Longltudinal 871 (1600) 979 (142 862 (125) 19.5 23.3

: VF-291

• 688-IOL-I Longitudinal 982 (1800) 676 (98 607 (88) 15.1 23.1

675-8L-2 _Longitudinal 982 (1800) 717 (104 641 (93) 11.7 __15"5

•Tensile tests were conducted uslng 1.8 mm (0.070 inch) diameter test specimens machined

from SC turbine blades that have been heat treated and then ex_sed to g82°C (1800°F) for

i000 hours. Prior to the 982°C (1800°F) exposure, heat treatment consisted of 1288"C

• (2350"F) for 2 hours, plus 1296eC (2365"F) for 4 hours with rapid Argon Quench, plus 982"C

(1800°F) for 5 hours, plus 871°C (1600°F) for 20 hours.

• *Longitudinal and transverse orientations are wlthin 15 ° of [001] and 15 ° of [i00], respec-i

tlvely. J

Stress versus Larson-Miller parameter relationships for times

' to one- and two-percent c_eep strain and rupture life a_e provided

in Figures 43, 44, and 45, respectively. Inspection of these

figures indicates a relatively large amount of scatter in one- ,_

percent-creep data. This extent of scatter is partially attributable _o mounting the extensometry on the grips (due to the small

size of the MFB minibars).

!

Isothermal stress versus rupture life data for SC NASAIR i00

and SC Alloy 3 are provided in Figure 46. These data indicate that

the rupture life curves at 871 °, 982°, and I093OC (1600 ° , 1800 °, .,6

and 2000"F) have the same slopes for both alloys. Consequently,

1 the SC NASAIR i00 alloy stress-rupture capability was essentially _

unaffected by the precipitation of the mu phase.

' 103 '_
ORIGINAL ?/_CEIS

OF POOR QUALITY

TABLE 47. CREEP-RUPTURE RESULTS ON 1.8 rnM (0.070 INCH) DIAMETER SPECIMENS

MACHINED FROH SC NASAIR i00 TURBINE BLADES.

I I o, d I**°I Ituptu_e 1.0 2.0

..°,.n .... .._.,... ,--.,, _.. "lon,n.on.o, • ,.°..r ,.*°.n*

Number Heat Orientation* NPn (kni) (hours) (t) (_). (hour .m.) (hour8)

648-4[.-1 Vr-235 Longl_udinnl 740 414003 842 (1253 192.4 12.8 18.1 20 55

S91-111..1 VW-0J5 LongL tudinnl 827 4120) 439.? 15.9 1D.1 84 159

148-3T-1 Vr-2$S Transverse 842 (123) li0.4 9.5 11.8 19 56

241-4T-1 Vr-235 Transverse 827 41203 214.3 11.2 15.3 11 41

$91-85.-1 VP-235 Transverse 7J3 4115) 354.6 10.3 14.2 32 116

511-121-1 VF-220 1`ongLtudLnnl 783 41153 444.9 12.3 14.4 12.0 184.5

4|1-91` V,-154 1`ongL todtnnl 758+793 4110+119"*) 790.2 8.9 13.2 ....

491-121` VF-1S4 1`ongL tudLnnL 758+792 4110 7.1 15.4 .... 

481-5T VP-IS4 Trnnnverie 758+793 (110,113"*) 551.9 10.8 23.2 ....

481-12T VF-154 Trnnlvetnn i 758 41103 308.8 3.e 10.0 ....

511-?T-1 vf-030 TcanJvn/no 871 (1600) 482 (70) 105.8 11.8 ;83.0 ;t;t.0 4i.0

509-41`-1 VF-154 1`ongi tudinnl 403 (703 144.4 18.1 24.? 27 55

840-15.-1 Vr-235 5.rnnmve[ne 393 (57) 330.51 12.8 17.9 104 103

848-31-1 VIP-235 1`ongL t ud_nnl 379 (55) 508.8 24,5 32.1 148 242

$11-111`-11 Vr-2;t0 1`on9 L Tudinnl 379 ($5) 458.0 ;80.2 27.5 01.7 178.5

848-75'-1 Vlr-235 5.rannver se 359 (52) 355.0 18.4 23.4 S0 158

$91-11T-1 vr-235 Trnnnver =e 345 (S0) 477.0 24.9 27.3 188 348

513-75.-1 VF°220 Transverse 345 (50) 519.0 18.1 25.3 74,0 211.5

140-81`-1 vr-235 1`onvt nod f,nnl tI 345 450) 844,0 ;82,5 34.3 167 305

. 581-191-1 V7-235 1`ongit udLnnl [ 331 (48) 1008.3 ;80.9 27._ 234 421 '_

591*;82T--1 V1'-235 Tcenevecae [ 3284 4673 050.J 16.0 17.) 164 3428 "' 591-221.-1 vr-235 1`ongl tud Lnal 310 (45) 1382.1 ;86.9 31.9 210 490

;. 511-10T-1 vr*220 TcnnnwerJe 902 (iq00) ;841 (35) .81. ;t 15.1 22.2 1.? 5.0

648-;tL-1 Vr-;t35 Longttudinal 241 (35) 125,8 25.9 38.7 -- 228

758-11°1 Vr-;t2? LongltudLnnl 241 (35) 64.7 12.1 28.;t ....

786-11`-1 vr-;t37 Longltudinnl J ;841 ()S) 86.2 13.2 31.7 ....

"" 782-11-1 W-237 1`ongi t_d inn1 I ;841 (35) 88.4 19.0 28.8 ....

! 458-31` VW-154 Longitudinal 207 (30) 188.1 24.. _ 42.6 ....

t 458-41` VF-XS4 Longitudtnnl ;t07 (30) 190.1 15.7 41.3 ....

481-3T V1'-154 5.[nnnunr se ;t0? (30) 174.4 15.6 28.3 ....

401-7T Vf-154 TrnnlverJe 007 (30) 189.5 i 26,0 34.;t ....

511*12T-1 Ylr*;t0o TcanlverJe 207 C]0) 209.5 i z0.0 83.s 9.2 31.1

848-;85.-1 Vlt-33S Tl_nnivnc,an 193 (28) 282.7 i ;82,8 29,6 -- 50

509-41`-2 VF-154 1`ongLtudinnl 179 (26) 783.0 ;88.3 31.9 ;87 239

513-91`-1 vr-;t20 Lon itudtnnl 104 (;873 482.;t 13.4 ;80.4 ;82.5 188.0

511-101`-1 vlr-220 Longi _udLnnl 1728 (;85) 818.5 20.7 ;87.9 30.5 303.5

513-9T-1 vr-22o 5.ranjvirln 172 (;853 427.9 11.1 15.3 69,5 497.3

591-15T-1, Vf-235 Trannverae 159 (23) 888.3 i ?.0 14.6 08 119

427-41`-2 Vr-154 Longitud_nnl 159 (23) 1(04.1 I 12.8 20.7 30 ;824

457-11`-1 VF-154 Tongltudinnl IS;t 4223 1143.9 19.6 31.8 40 729

511-91-1 VF-2220 Longitudinal 1038 (19003 131 (19) 451.0 7.2 14.5 12.0 331.5

513071-1 vr-;t20 Long_ tudinnl 1068 419503 131 (193 249.1 ! 10.? ;iS.3 124.3 233.8

511-9T-1 VF-2220 ?ranlvtrin 1G93 (2000) 131 (19) 89.2 5.3 10.1 9.2 56.7

509-91-1 VF-154 Long/tudinnl 131 (19) 90.7 128.5 ;81.7 -- 33

S09-iL*;t vr-lS4 LongLtudinil 110 414) 383.4 10.2 21.0 ....

458-81` VF-154 1`ongi tudinnl 103 4153 501.0 11.2 34.5 ....

459-111 V!'-154 LongL tudinnl 103 (15) 453,? 10.8 33.7 ....

511-11T-1 Vr-;t20 Tranaverle 103 (15) 123.4 5.3 8.1 17.3 2214.7

I

511-71`-1 uP-;t20 1`ongitud_nnl 103 (153 273.1 0.4 30.0 19.5 ;it0.1

t 513-61-1 VF-220 LunRudt nil 90 (13_ 380.8 11.2 228.0 11.2 319.3

513-8T-1 VF-;t_0 Tcnnlverle 90 (13) 335.6 1.9 12.3 ....

868-8T-1 Vr-235 TrnnlVnraet 83 412) 810.5 4.;t 0,3 10;t 400

591-211-1 VP-235 Longitud inn1 ?i (113 1234.S 9.4 ;iS.8 611 1057

4281-41-1 VF*154 Lon ltudt nnl 89 (L0) 2055.;t : 14.6 17.9 ....

591-;83T-1 V1_-235 Trinsvnrsn 49 (10) 1707,9 9,7 17.2 -- 471

/

458-9L01 Vr-154 Longi Tudlnal t 628 (93 3188.0 lm 10.? 13.1 500 11289

Hint trnntoan_ oonlllted of 1318*¢ (;t400*r) for hOUri, plua 1324*C (;84150P) {Or ;t hour| wz_h caped &coon Quench.

plul 913"¢ 41000*r) /or S houri, plui (171'C (IiOOelP) for 20 houri.

tpnH0iiln Ortlflti_ofli lIFt vLthtn 13 drrill O{ cOOt)*

• *|pecilinl vice uploaded to 793 14Pi 1115 kli) niter S00 houri.

104
_' ORIGINAL _'"__z,,.,'.;,_" i'%,

OF POOR QUA,__'rY

,f.,

1379 12001

; BEST FIT

"_, ego11oo_ _ _,_ ,

-, 414 (60, "l

,: -- 276 1401 MINUS THr_EE SIGM ., ,

T_ IO

._ BL

,"' _p. 138 (20) & ,

% ,

69 1101 _

55 :B) •

41 161 ,'

.. 34 37 40 43 48 49 52 55 58 ._

F = 1.8 (T + 273)(20 + LOG T) x 10-3

P,

._2 Figure 43. 1.0-Percent Creep of SC NASAIR i00, 1.8 mm

/ (0.070 Inch) Round Specimens.

, _

1379 12001

f _ BEST FIT

690 11001

'_' = 414 1601

p, r,n

m _MINUS THREE S o. 276 1401

n-

• t- 138 1201 VJ

69 1101 .__ _ _, '

55 181

41 (0) -,,

36 39 42 45 48 51 84 57 60

b I

P," 1.8 (T + 2731120 + LOG T) x 10.3 t

I

Figure 44. 2.0-Percent Creep of SC NASAIR I00, 1.8 mm

(0.070 Inc:_; Round Specimens.

105

I

i

*
OF POOR QUALITY

l_g (20o)

6110(100) _'='_ ./s_,'r,=,T

,_1 414 (60)

- f __

1_69 120 |10) ) " _ '_"_,.. 68(8)

41 (6) .......

37 40 43 46 4A 52 55 58 61

P,. 1.8 (T + 273)(20 + LOG T) x 10-3

Figure 45. _tress-Rupture of SC HASAIR 100, 1.8 mm

(0.070 Inch) Round Specimens.

Microstructural examination of SC HASAIR i00 specimens evaluated at 982 ° and 1093°C (1800 ° and 2000°F) with rupture lives less

than i00 hours and the longest rupture lives at 760 °, 871 °, 982 ° ,

an(] I093°C (1400 °, 1600 °, 1800 °, and 2000OF) are provided in

Figures 47 and 48 respectively.

Microstructural examination of the SC NASAIR I00 microstructures after creep-rupture testing indicated the following:

. o At 760°C (14000F) and 814 MPa (118 ksi), the microstruc-

* ture remained stable during the 563.8-hour test. This

result is consistent with alloy development results,

i which indicated that the microstructure remained stable

after a 790-hour exposure at 760eC (1400°F) [500 hours/

759 MPa (ii0 ksi) + 290 hours/?94 MPa (115 ksi)].

106

}
ORIGINAL P:%C_ I_

OF POOR QUALITY

NASAh100 ; ALLOY3

13791200)- -

137912o01 -[-/ 760oCi1400°Fi 760oC(!400° Fi

69011001-L- e"_'i_-_'_e'_ 690(1001-_ o o

_ 871°CII600°F) _._ _ 671°C[1600F°_

_-- - _...._ _82° C(IBOO°FI _- _ C(i.800oFI

X - ._.-.o-. _ L,.._ E -

;.'i W'J - 0"43"--.,-....

P' "'__"'- _ _ 60[10) _"_ -- '_l'-.--...l,2 IO03°C'2000°F'

,_ 69110l- -_

3a(51IL___ l I I t I I i I{ _ I I { 34(51 I ,Jl{ _L_._L_L_ al, 'i ' _" :,

50 IO0 500 I0110 5000 50 !00 500 lOoO 5000 ,_

: RUPTURELIFE.HOURS RUPTURELIFE,HOURS

"C

i-'. cz_ooUNCOATEO

"'- ai,_O ALUf,,IINIDE COATING

2; _ .

,_,_ Figure 46. Stress-Rupture Life of SC Alloy. _. ;,2.

o MU phase precipitation was observed in short-life speci- "

mens (less than i00 hours) tested at 871 °, 982 ° and

I093°C (1600 ° f 1800 ° l and 2000°F) • This result indicates _j

that the mu phase was present in the long-life specimens _:_

., for greater than 93 percent of the rupture life and did _

not appreciably affect it. ",

O Propensity for mu phase precipitation was greater at ",.,,

982°C (1800°F) and I093°C (2000"F) than at 871°C "

" (1600"F) . <

_t

o All SC NASAIR i00 and SC Alloy 3 creep-rupture-tested M_'B

> specimens that were metallographically examined during ,_

this task revealed that secondary creep cracks were com- _-

_,. monly associated with microporosity. . :

107 _4
ORIGINALPAGE iS

OF POOR QUALITY

_ .-, ., - p_'". ],_.j_: {_,/ " ..

• ,'"_"_." ." "/ - ._

t

o . .

..¢.._ .... _ _"- -. _"

982°C (1800°F)/241 MPa(35 KSI)/86.2 HOURS .. _

-- ,.. ..:. ,, _;:,:_

.-" . I•- ,'-. - 4

. " . ,.._....._.. :.

• - .-' .

. ". ,.,.- ! -< "'._ ._._

• "'=- "'_L " _.._ _;

, ,_:-,.....,.:_, . .-_ . , ..

.. _._,,_ "" '. ".' .... I '_. _'.'_._ "

:...._..-.....__,._,_...:.,__. .-_-_ ...._,:.:,._.:__,

E ._,',;-_,-._"_"-_"*;: ' ,__iK"_._,__-_"_'_'_'_-__

P772474

1093°C (2000°F)/131 MPa (19 KSl)/69.2 HOURS

Figure 47. .Short-Life Microstructures of Creep-Rupture-Tested

SC NASAIR I00 MFB Specimens. 400X Magnificatio[_.

.. 108

,]
ORIGINALPAGE _S

OF POOR QUALITY

: t"

t w

• ! I

I

/,

L \

-_,_,.', _ ...:,_. • ........................... • ...........

'_ (lOOX) 1400X)

760°C 11400°F1/814MPa (118 KSl)/563.8HOURS

4-."

_ _._p_,,, ,

,;, t ' "- "

• i:

1

• , . • 4.- .a . _P"

r2 " _

i

•_ P77247-4

'_ (100X) 1400X) "

', 871°C(1600°F1/310aPa (45 KSI1/1382.1HOURS i

t Figure 48, Part I. Long-Life Microstructures of Long-Time

Creep-Rupture-Tested SC NASAIR 100 MFB

,, Specimens , .

Y 109

I'
ii0
ORIGINAL PAGE iS

OF POOR QUALITY

For some superalloy systems, application of protective oxidation resistant coatings can severely degrade the creep-rupture

life. 14 Consequently, creep-rupture tests were performed with diffusion-aluminide (RT-21) coated SC NASAIR 100 specimens in order to

evaluate potential coating effects. Results obtained in this evaluation are provided in Table 48. Two specimen stress levels are

shown in Table 48, one that includes the coating (specimen stress)

and a calculated stress that uses an area consisting of the base

material minus the coating (substrate stress); the substrata stress

l

assumes that the coating does not support any significant load.

Orientation of all coated test specimens is longitudinal and within

15 degrees of [001]. ._.

Computer analysis of the coated creep (substrate stress data)

- versus uncoated data revealed no statistical difference in the time

•- to 0.5-, 1.0-, and 2.0-percent strain or rupture life indicating no

"J degradation of SC NASAIR 100 creep properties _com application of a

•- the diffusion aluminide RT-21 coating. The coated data are also _

compaced with the uncoated data in isothermal stress versus rupture

life curves (Figure 46) to illustrate the above conc]usion.

a&

TABLE 48. CREEP-RUPTURE TEST RESULTS ON DIFFUSION-ALUMINIDE (RT-21) COATED I.SMM

(0.070 INCH) DIAMETER TEST SPECIMENS MACHINED FROM SC NASAIR i00 BLADES _ • _/

..... ,..... , ................. %

7 ISDecl:men ISubstrata 1Rupture 1 I ,0._5 1.0 ,2._0 ,

Mastel___.._j j

S_)eclmell J LIf J _ell_rltuEe Stresm Stres_ Lira i Elongation R of A Percent P@rcent Percent Humber J_Xe O_zent__tzon" "C I'F) [NP_a lksz) NIPS (kill) lHoursl J (ll) J (I) (aouzs_ (Hou,s[_ (_Hflo_urs._.__) ,

SlX-eL-z ivt-no [Long[tudznal 760 1*4oo1 152 11o911el4 ISle1 ! S_],e I e,s I 1s.9 n,2 zad.e _6o.o

511-6L-1 I[VF-220I{Lonqitudindl 871 (Z4001 ]S9 152)I 3a6 (54) I s3J.9 I 16.e I 19.0 2e.1 lie.2 22].0 'j

I I

513-4L-1 I W1'-220 I 345 150) I 770.2 I 16.1 I 21.3 6B.5 213.0 370.7 ' _ ILongl tudlnil

871 (16001 117 (461

; SII-SL-] Ivr-220 g8_ 1|800) 172 (2s) 1186 12?1 I 478.4 I 16.1 I 23.5 21.5 82.e 335.9

,ongJ cudlnel ,

513-3L-! I VP-220 Longit ,dine1 982 (1800) 152 1221 ] 172 12S) I 12S.] I 16.6 I _2.T _S.0 21_.4 _93.6 ,_,,_.

,4:, 521-41,-1 I VF-220 ILong*r.udz,al 1038 (1900) 124 (111) ] 138 (20) ( 759.3 | 13.9 I 21.4 0.3 ...... J _r

I '

513-2L-1 IVr-220 [.ongitud*..; ' IG?e 119o01 97 1141 II0) 11_1 I1i14.4 I ll.] I 20,2 7.§ 351.1 1354.4 I _

• , 51|-ZL-1 Ivr-2201Longl'_dln*l 1093 (20001 83 11_1 I _0 (1_1 I 44 .3 I l).0 I 21.7 I 1.4 9.1 I Ze0.O _

I

,*_ i

,', [ l 1 lo,(,ooo, i ,,,,11I,o,, ""1'"°, I ,o.l11,.,E,., _..

_' Heat treltlant co_l|eted o[ 1316"C 12400"r) to¢ 2 houri l plUl 1324"C (211S'r) for 2 houri wllh

rep_d Argon (_usnch, plus RT-21 coated, plus 671"C (1_00"1 _) lol 20 hours.

•8ps(.lmen oriente_ione Ilia all v/thin IS dsgcees o_ (0Ol) direct|on.
ORIGINAL PAGE iS

OF POOR QUALITY

Creep-rupture tests were also performed to assess the effects

of a 982°C(18000F)/I000 hour exposure on the rupture life. As previously indicated in Figure 40, this exposure coarsens the gamma

prime phase and permits the mu phase to precipitate in NASAIR i00.

Stress-rupture results for exposed SC NASAIR i00 are compared with

similar unexposed data and with data for exposed and unexposed SC

Alloy 3 in Table 49 and Figure 49.

Examination of these data reveals the following:

o SC NASAIR i00 and SC Alloy 3 data obtained after the

• 982°C (1800°F)/1000-hour exposure was contained (with

two exceptions--one high and one low) within the typical

and minus three sigma lines for fully processed NASAIR

.% I00 •

?

O After the 982°C (1800°F)/1000-hours exposure, typical

values for both SC NASAIR i00 and SC Alloy 3 appear to be

approximately one-half a Larson-Miller parameter lower

than fully processed SC NASAIR I00.

_ TABLE 49. RUPTURE LIFE COMPARISON OF SC NASAIR i00 AND

: ALIOY 3 BEFORE AND AFTER EXPOSURE AT 982°C

; (1800"F) FOR i000 HOURS.

d

' Rupture Life

(Hours)

After

Exposure

Test Fully at 982"C

-i Temperature Stress Heat (I800OF) for

Alloy "C ('F] MPa (.Ksi) Treated i000 Hours*

NASAIR 100 982 (1800) 241 (35) 125.8, 81.2 100.6

Alloy 3 I 241 (35) --- 76.8

r

HASAZR i00 i 207 (30) 189.3, 174.4 95.1

NASAIR i00 198.1, 189.5 196.8

.^SAXlO. O 2o9., -

Ahoy3 i 1,0.2 160"3

Alloy 3 161.3 191.7

NASAIR i00 186 (27) 482.2 268.0

Alloy 3 186 (27) 278.2 208.4

NASAIR i00 i 172 (25) 616.5, 627.9 200.8

Alloy 3 t 172 (25) --- 484.8

, *Fully heat treated prior to exposure.

112

®
ORIGINAL PAGE E._

OF POOR QUALITy

345(00)

276 (40) _ _ NASAIR 100

•.: =c4 207 (30) -' / -_ - -

m NASAIR 100

(FULLY PROCESSED)

: m" BEST FIT MINUS

THREE SIGMA

,0.

! ;

(]) ALLOY 3 (EXPOSED

,t

¢,

z"

69 (10) _

:" 47 48 49 00 51 52 53 54 L_

P- 1.8 (T + 273)(20 + LOG T) x 10-3 I

Figure 49. Exposed - 982°C (18N0°F)/1000 Hours - SC NASAIR i00 )

and SC All.oy 3 Rupture Data Compared to SC NASAIR

. i00 Larson-Miller Stress-Rupture Curve (Fully

Processed ). ,_

" These reductions are thought to be predominantly associated "

with gamma prime coarsening. To support this conclusion, life •

l

i reductions were comparable to those observed when the SC NASAIR i00

i microstructure was degraded by gamma prime coarsening as a result .

1 o_ improper (slow) cooling from the solutioning temperature (com- ',

i '

pare Table 30 with Table 49). i_

113 ii
SC Alloy 3

Comparative creep-rupture data for SC Alloy 3 MFB minibar

specimens were obtained for the same conditions evaluated for SC

NASAIR 100. Uncoated and coated creep-rupture data obtained in the

program are provided in Tables 50 and 51, respectively. As with SC

NASAIR 100, analysis of the data indicated that there was no significant difference for longitudinal and transverse specimens as long

as the specimen orientation was within 15 degrees of an <001>

direction.

Stress versus Larson-Miller parameter relationships for times

to one- and two-percent creep strain and rupture life are provided

in Figures 50, 51, and 52, respectively. Coated and uncoated data

are provided in these figures and demonstrate (on a substrate

stress basis) that the RT-21 diffusion aluminide coating did not

affect the creep properties.

Isothermal stress versus rupture life data for SC Alloy 3 (and

SC NASAIR 100) are provided in Figure 46. Examination of the data

in this figure and the data provided in Tables 47 and 50 indicate

that SC Alloy 3 is slightly weaker than SC NASAIR 100 over the 138

to 793 MPa (20 to 115 ksi) stress/760 ° to 982°C (1400 ° to 1800"F)

temperature ranges, which are of primary importance to turbine

blade design.

Post-test microstructural examination of SC Alloy 3 specimens

revealed no evidence of TCP phase formation after 1685.5 hours at

982°C/152 MPa (1800°F/22 ksi) and 2835.5 hours at i093°C/76 MPa

(2000°F/II ksi), indicating elevated temperature phase stability.

Microstructures of longest rupture life specimens at 760 °, 871 °,

982 °,and I093"C (1400 °, 1600 °, 1800o, and 2000°F) are provided in

Figure 53.

114
/

ORIGINAL P_.GE;,_

OF POOR QUALITY

TABLE 50. CREEP-RUPTURE RESULTS OH 1.8 MM (0.070 INCH) DIAMETER SPECIMENS

MACHINED FROM SC ALLOY 3 TURBINE BLADES.

Tl_ to Tine to

lupturn l.O 2.0

Jpeclnan Haste: Teel_er_curn Stress Lifo Elongation J of A Percent Percent

Number Host Ori_ntation* C (W) NPa (knt) Shouts) (0) (tp (hours) (hoursl

;. 600-1L-1 VF-lgl Longitudinal 760 (1400) 793 (1152 106.5 IS.I 19.5 3 10

675-1L-1 vr-2$1 Longitudinal 760 (145b) 758 (IlO) 294.3 14.0 20.6 6 19

,_ 483-1L VF-99 Longitudinal 760 (1400) 758 (110) $170| 15.4 23.7 ....

483-2L VF-99 f,cmgltudlnal 760 (1400) 758 (110) 291.8 16.% 2_.3 ....

? 684-2L-1 Vr-2_l f,c_gitudlnsl 071 (1600) 379 (55) 378.0 11.2 14.6 14 126

688-1T-1 VW-291 Transverse 871 (1600) 365 (53) 459.2 11.2 20.4 151 248

688-2T-1 VF-291 Transverse 171 (1600) 345 (50) 712.5 15.1 23.7 125 323 ._

" 675-3L-1 VP-191 Longitudinal 871 (1400) 310 (45) 1052.0 15.4 27.4 193 425

792-1b-1 VF-292 [_ngltudlnel 982 (11002 241 (35) 51.6 21.7 29.8 ....

794-1L-1 VF-44_ Longitudinal 902 (1000) 241 (35) 70.4 20.3 2T.3 .... :

483-2L VF-gg L_ngltudinsl 982 (18002 207 (30) 180.2 22.5 28.9 .... _

_" 483-0L VF-99 Longitudinal 902 (1800) 207 (302 141.3 10.2 20.2 .... :

675-4L-1 vr-191 [,ongltudlnsl 902 (10002 Lid (27) _70.2 17.4 29.8 30 125

!- 684-3b-1 vr-291 Longitudinal 982 (18002 152 (22) 1335.4 12.8 19.1 80 5t4

4l]-9L VF-99 LOngitudinal 10J3 (2000) 163 (15) 247.0 11.0 20.0 ....

483-10L VF-99 _ongLtudlnal 1093 (2000) 103 (15) 199.1 16.g 31.0 ....

, 475-5T-1 Vr-29_ Transverse 1093 (2000) 103 (15) 320.0 5.3 12.1 -- 48 :

,* 675-5L-1 vr-291 Longitudinal 1093 (2000) 90 (13) 581.9 7.5 13.1 -- 354

_- 675-6T-1 Vt-291 Transverse 1003 (2000) 83 (122 1431.1 5.0 6.9 -- 711

618-4L-1 Vr-291 Longitudinal 1093 (2000) 76 (112 2035.5 9.4 11.7 530 1090

*Npeelman o_Lantatiou8 mire vithin 15 degrees of <001>.

Neat trearnent consisted of 120Nee (23504F) for 2 hours, plus 1291°C (23654F) tar 4 hours

vith rapid Argon Quench, plus 982eC (I800eF) for 5 hours, plus 0710C (1600oF) fo_ 20 ho_rs.

TABLE 51. CREEP-RUPTURE RESULTS OH DIFFUSIOH-ALUMIHIDE (RT-21) COATED 1.8 MM "

(0.070 INCH) DIAMETER TEST SPECIMENS MACHINED FROM SC ALLOY 3 TURBINE

BLADES.

Ties to Tins to

Specimen Substrata Rupture 1.0 2.0

Specimen 14_stsr Temperature 8trese 8Crns8 Lifo Elongation R of A Percent Pe:uent

Number Heat Orientation *C ._er) NPa (ks1) NPa (kel) (hours) (t) (t) (nours) (hours) t

675-2_-L _r-191 Longitudinal 942 (l|o0p 152 1211 172 1251 73_.7 11.6 20.0 23 102 ,

j 688-5:.1 vff-251 Longitudinal _02 (18002 138 (20) 152 (22) 1605.5 11.3 21.1 210 l)40 ,

_es-6_-I v]r-2_l LongL_udLnel 1093 (2000) 83 (1_) 90 (13) 007.4 ).e 11.2 - - ,_

4ii-lL-I Vf-191 Longitudlnsl 109J _000) 69 (10) 76 (11) 1532.6 7.4 l_.l 190 700 _

Heat trsstment nonnlnted of 1280*C 11350"_1 for 2 hours, plum 12SSeC 11165*r1 for 4 hour8

vith rapid Argon Ouen_h_ plus n_-21 coating, plus 871"C (16000r) foc _0 hours.

115
ORIGINAL PACE iS

OF POOR QUALITY

1379(200) .........

x2

_:414181 M,Nu., T.._ESlGMA__ \

_27! (401

Q:

.." 13812B) ,,

UNCOATED

Jk RT-21 COATED , ,:

IO(10).,. '

37 40 43 46 49 S2 55 t58

P - 1.8 (T + LOGT) x 10.3

Figure 50. 1.0-Percent Creep of SC Alloy 3, 1.8 mm (0.070 Inch)

Round Specimens.

6.2.3 Low-Cycle Fatigue

Axial strain controlled low-cycle-fatigue (LCF) tests were

performed at Mar-Test, Inc. of Cincinnati, Ohio. These tests were

performed using 6.35 mm (0.25 inch) gauge diameter specimens (Figc ure 54) machined from cast slabs (SC NASAIR i00) or bars (SC Alloy

_, 3). Data were obtained at 760°C (1400"F) at a 0.33 Hertz (20 cpm)

i. frequency and an A ratio (AlternatingMeanStress StressAmplitude ) of infinI

ity. Test results are reviewed in the following paragraphs.

T

i 11, ; I

kr
ORIGINAL PA,_,_ iS

OF POOR QUALrI'Y

1370(200),

690 1100}

,. 552(80, ___

" _,_%_BEST FIT

414 (60)

• _-. MIN

276 1401 •

138120)

_+" A UNCOATED

& RT-21 COATED

+

"_: 69 (10) i1'

36 39 42 45 48 51 54 57 i

, P = 1.8 IT+ 2731(20 + LOG T) x lr 3 _,

t

Figure 51. 2.0-Percent Creep of SC Alloy 3, 1.8 mm

(0.070 Incn) Round Specimens.

SC NASAIR I00

LCF tests of SC NASAIR 100 weJ:e conducted at total strain

ranges between 1.2 and 2.5 percent. Test data are provided in

1

Table 52. Computer generated best-fit and minus-three-sigma strain

i,

range versus cycles to fatigue curves are provided in Figure 55 for

'_ longitudinally oriented specimens. The average LCF curve indicated i ,.a

that a total percent strain range of 1.45 percent would yield an

LCF life of i0,000 cycles for longitudinally (within 15° of [001J) I

oriented SC NASAIR i00. !

,i

'+ 117

' ,I(i',,
ORIGINAL PAGE 18

OF POOR QUALITY

1379(200) ....

000(100) _'_ •

"x',.

6621001 "_ nESTFIT OF ACTUALDATA

.l

• 414(60)

s/

1381201

UNCOATEO

• RT-21COATED

00IlO)

30 42 46 U 61 64 67 60

P -'-1.8 (T* 273)(20 + LOGT) x 10-3

Figure 52. Stress-Rupture of SC Alloy 3, 1.8 mm

(0.070 Inch) Round Specimens.

Macroexamination at 40x and SEM examination at higher magnifications were performed on the fracture surfaces of the LCF tested

specimens. In all instances, cracking in these specimens initiated

at microporosity (Figure 56). Cracks typically propagated initially in Stage II, approximately perpendicular to the applied

_'" stress in the (001) plane, and then switched to Stage I and propa-

" t

_ , gated on a crystallographic plane. The duration of the initial r

Stage II mode of crack propagation seemed to be greater for longer

life specimens that were tested at lower stresses.

, i18
ORIGINAL _"P¢ 1IFt-*' " _d L,, [_

OF. POOR QUALITY

• i

", qlp

f

, I

i i -4,4'

7.

; 760°C (1400°F)/758 MPa (110 KSI) 871°C (1600°F)/310 MPa (45 KSI)

; 294.3 HOURS 1052.9 HOURS

' _ ...... o _-." ._

q.,._, - _ -_, -_

"-.<

,_,

982°C.................. (1800°F)/152 MPa (22 KSI) 1093°C (20000F)/76 MPa (11 KSI) P77247-7 0.","'!

1685.5 HOURS 2835.5 HOURS "_

_' i Figure 53. Long-Life Microstructures of Long-Time Creep- ""

i Rupture-Tested SC Alloy 3 MFB Specimens _

_L (400X Magnif icat ion) . 119 ¢'i"
ORIGINALPAGE iS

OF POORQU/-tLITY

63.5mm

i (2.50 INCH)

12.7 mm

(0.5 INCH)

25.4 mm _,, 12.7 mm DIAMETER

(1.0 INCH) _ 1-20 UNF3A THREAI)

RADIUS 6.35mm 2

-- (0.25 INCH)

DIAMETER

Figure 54. LCF Test Specimen.

TABLE 52. LOW-CYCLE-FATIGUE TEST RESULTS - AXIAL-AXIAL STRAIN-CONTROLLED TESTS,

6.35 _ (0.25 INCH) DIAMETER TEST SPECIMENS

MACHINED FROM NASAIR i00 SLAB CASTINGS.

Total

Strain Elastic Modulus

Specimen A Temperature Range NE

Number Orientation* Ratio OC (°F} (t] 106 GPa (I06 psi) (Cycles) Remarks

.... ...... i___i

II _Longitudinal _ 760 (1400) 2.50 84.A (12.2) 427

1 Longitudinal _ 2.40 78.6 (II.4) 750

12 Longitudinal _ 2.20 82.1 (ll.9) 2,005

3 Longitudinal _ 2.00 80.7 (11.7) 2,677

8 Longitudinal _ l. H0 84.1 (12.2) 2,133

5 Longitudinal _ 1.70 83.4 (12.1) 4,283

7 !Longitudinal _ 1.60 83.4 (12.1) 5,450

2 lLongitudz-al = 1.50 82.1 (11.9) 15,845

6 Longitudinal _ 1.40 80.O (11.6) 14,540

I0 Longitudinal _ i 1.30 85.5 (12.4) 15,060

9 Longitudinal _ 1.20 77.9 _11.3) 29,600+ Teat terminated

16 Transverse _ 1.50 80.7 (11.7) 1,816 Shrinkage porosity

i_eat treatment consisted of 1316"C (2400"F) Eor 2 hours, plus 1324°C (2415eF) for 2 hours

with Rapid Argon Quench, plus 982°C (1800"P) EeL 5 hours, plus 871°C (1600°F _ fnz 20 hours.

A ratio - alternatlnmean o stress amplitude stress , Frequency • 0.33 Hertz.

Longitudinal within 15 degrees of [001] and Transverse within 15 degrees of (100]

120
ORIGIN_L _'__"-

OF POOR QU_LFi?

ACTUAL DATA

I-Z

°

: .J THREE SIGMA

• z

<

: _ 1.50 -' -_

_z

¢¢1

.J

< 1."nw .....

O ;

i- _,TEST TERMINATED

/ 0.50

"- 100 1000 10,000 10(1,000

Nf (CYCLES TO FAILURE)

Figure 55. Strain-Controlled 760@C (1400OF) LCF _est Results _

; for Longitudinal 6.35 mm (0.25 Inch) Diameter

Test Specimens Machined from SC NASAIR 1G0 Slab i

Castings.

t

The single trans-ersely oriented NASAIR i00 specimen that was

L_ _ tested at 1.50 total percent strain range failed in 1,818

cycles, which was significantly lower than the longitudinally

oriented specimen tested at the same strain range. SEM examination

of the transversely oriented specimen fracture surfaces revealed

substantial (unacceptable) shrinkage porosity, which explained the

lower life.

SC Alloy 3

%

,' LCF data for SC Alloy 3 were acquired at total strain ranges

between 1.2 and 1.8 percent, tePt data are provided in Table 53.

q

{

121

;
ORIGINAL PAGE IS

OF POOR QUALITY

760=C (1400°F)/1.30-PERCENT RANGE/A =0o/0.33 Hz

• , ,,,¢'r?_''r'_'' *-'<_

,,:

3 _-J" . .=/

, /./' .-,,.;,;-... ,, ;

.¢ ',_b w

- . ..... ww==

200X lOX .":-.

INITIATION SITE: CASTING POROSITY INITIAL PROPAGATION: STAGE II

FINAL PROPAGATION: STAGE I

(LIFE = 15,060 CYCLES)

Figure 56. LCF Fractu_:e of SC NASAIR i00.

A comparison of the SC Alloy 3 data with that for SC NASAIR

i00 is provided in Figure 57 on a strain range basis. This comparisop indicates that SC Alloy 3 has a lower cyclic strain ranje

capability than SC NASAIR 100.

The elastic modulus of SC Alloy 3 was about 20 percent greater

than that for SC NASAIR 100. Consequently, on a stress range

basis, SC Alloy 3 and SC NASAIR 100 are essentially equivalent (as

shown in Figure 58).

122

...... i
ORIG!NAL P,-,_,- !._

OF POOR QUALITY

TABLE 53. LOW-CYCLE-FATIGUE TEST RESULTS - AXIAL-AXIAL STRAIN-CONTROLLED TESTS,

- 6.35 MM (0.25 INCH) DIAMETER TEST SPECIMENS MACHINED FROM SC ALLOY 3

SEPARATELY CAST BARS.

Total

Strain Elastic

Specimen A Temperature Range Modulus Hf

Numbez Orientation* Ratio "C (eF) (Percent} 106 GPa (106 p_i) (Cycles)

687-5 Longitudinal = 760 (1400) 1.8 ...... **

687-6 Longitudinal _ 1.8 97.9 (14.2) 756

687-4 Longitudinal _ i 1.6 101.4 (14.7) 1,510

I

692-1 Longitudinal = 1.4 97.9 (14.2} 6,071

687-7 Longitudinal = 1.3 96.5 (14.0) 9,668

687-3 Longitudinal = 1.2 100.0 (14.5) 14,545

Heat treatment consisted of 1288"C (2350"F) for 2 hours, plus 1296*C (2365"F)

for 4 hours wzth Rapid Argon Quench, plus 982*C (1800"F) for 5 hours, plus

871°C (1600"F) for 20 hours.

A ratio = alternstin_meastress n stressamplitude, Frequency - 0.33 Hz

*within 15" of [001]

**Equipment malfunction

>

2.50 :__

_ NASAIRFIT100OF ACTUAL DATA _:

THRSESI

_ 2.1111 _

! '

.

'I

O ALLOY 3 TEST REgULTS

t_

1,00 -' ,_

100 1000 10,000 100,000 ,,_

Nf (CYCLES TO FAILURE) _'

4

Figure 57. SC Alloy 3 Strain-Controlled LCF Test Results !_ Compared to SC NASAIR i00 LCF Curve at 7b0°C

(1400°F). _

123 J
ORIGIN/LL PAGE 19

OF POOR QUALITY

760°C(1400°F)/A = /0.33 HERTZ

4.0 --

o_ 3.0 -

_ 2.0 - NASAIR!00

..i _e'

I-- m ALLOY3 -'_ _'_._

I-,

1,0 J J I 1 I i Ii II l I I I I I II I; I I I I ! I I II i

2758(400i- - u.J

:m _ 206913001- ,,_,,,,_ASAIR 100

c_ _ 13791200)

¢/g a,.

OC

I,-

690(I00] ' .I i t I ,,'_, ! I I , _ ,,,,, I ..i I i i If,,, I

102 103 104 105

CYCLESTOFAILURE

Figure 58. Low-Cycle Fatigue of SC Alloy.

In subseque,. LCF tests performed under a Garrett R&D program,

the modulus of SC NASAIR i00 was found to he equivalent to that of

SC Alloy 3. Reasons for the relatively low modulus for SC NASAIR

i00 (Table 52) were not determined.

, Examination (40X) of fracture surfaces on the SC Alloy 3 spect imens LCF tested at 760°C (1400°F) revealed subsurface Stage ZI

r

initiation and Stage I crack propagation. SEM examination of the

fracture surfaces revealed that the initiation sites are associated

with microporosity. The fracture mode on the SC Alloy 3 specimens

, was similar to those observed cn SC NASAIR i00 specimens.

L 124
ORIG!NAL PAG_ tS

OF POOR QUALITY

' 6.2.4 High-Cycle Fatigue

Axial HCF tests were performed under load control by Metcut

Research Associates in Cincinnati, Ohio. These tests were performed with both smooth (Kt = 1.0) and notched (Kt = 3.0) 6.35 mm

(0.25 inch) gauge diameter specimens (shown in Figures 59 and 60,

respectively). Data were obtained at A ratios of infinity and

0.95, 871 ° and 982°C (1600 ° and 1800°F), at a frequency of 60 Hertz

(3600 cpm). The HCF data were obtained with specimens in the

uncoated, RT-2i diffusion aluminide coated, and exposed [982°C

(!800°F)/1000 hours] conditions. Test results for SC NASAIR 100

and SC Alloy 3 are reviewed in the following paragraphs.

A Ratio Effects

?

The effect of stress ratio, A = Alternating stress amplitude/

mean stress, on HCF life was evaluated with smooth specimens at ,

871°C (1600°F). Test data are provided in Tables 54 and 55 for SC

NASAIR 100 and SC Alloy 3, respectively.

,__F25.4 nm (1.0 INCH) RADIUS

6.35 mm (0.25 INCH) DIAMETER

-

,-_ 19 mm

_ (0.75 INCH)m '_5.g mm DIAMETER _v t

!_ 70 mm (2

_: i" 18NF-3ATHREAD

" Figare 59. Smooth HCF Test Specimen.

g2

125

)
ORiGiNaL PAGE i_

OF POOR QUALITY

m 15.9 mm (0.63 INCH) DIA. _ 32 mm (1.25 INCH) RADIUS /

5 18 NF - 3A THREAD 0. / m. (O.,,,N

_ _ 9.0 mm 10.35INCH) DIA.

llllf_ _ -_lllllllllL];

19(0.75mmINCH)m

70 mm (2.75 INCH) _--

4

: __ 0.33 mm (0.013INCH) RADIUS

Kt = 3.0

Figure 60. Notched HCF Test Specimen.

Test data are also compared graphically in Figure 61 with similar data for DS MAR-M 247 generated previously during Garrett's

, NASA-MATE Project i. 1 Inspection of this data indicates that the

107 cycle runout strengths of SC NASAIR I00 and SC Alloy 3 exceed

that of DS MAR-M 247, with SC NASAIR i00 having the highest

strength. As expected, pulsating tension conditions (A = 0.95)

|

resulted in higher maximum stress and lower alternating stress

amplitude at 107 cycle runout for each alloy relative to the cycle

with equal tension and compression (A = infinity).

Examination of the fracture surfaces of these specimens

revealed that cracks typically initiated at subsurface microporoslty. Cracks typically propagated initially in Stage II,

126

®
ORIGINAL F_GE i..S

OF POOR QUALITY
r %

ORIGINAL PAGE i_

OF POOR QUALITY

TABLE 55. EFFECT OF A RATIC OH HIGH-C_CLE-FATIGUE LIFE OF SMOOTH 6.35 MM (0.25 INCH)

DIAHETER UNCOATED SC ALLOY 3 SPECIMENS.

I I i -- .....

Stress 14We (ksi) Cycle:

I Spgiltn Nlster A Temperature ]ALternating to

Number J Beat Conficjuration Ratio J, *C (*F) Haxieu.m.J AmL)litude failure I. Remarks

I

. . .. ..

689-4 Vr-291 Smooth J co 871 (1600) 446 (65) 448 (65) 80,000

789-1 VF-29L Smooth | c.o 414 (60) 414 (60) 360,008

582-1 Vr-291 Snooth c_o 379 (55) 379 (55) 7,000 Unacceptable porosity

709-7 VIP-291 Smooth co 345 (58) 345 (58) 1,561,008

695-2 VP-291 Smooth _ 318 (45) 318 (45) 10,088,088 - test termLnsted 

502-1 VP-291 SmOoth c.o 3_-8 (45) 310 (45) L8,008,000+ Runouc - test tetnLnated

789-5 Vr-291 Smooth 0.95 483 (78) 234 (34) 415,808 PorosLty

£69-6 VP-2ql SnQoth 8.95 448 (65) 221 (32) 5w398,080

709-9 V]_-291 S8_oth 0.95 414 (60) 280 (29) 7,950,008

666-7 Vr-291 snooth 8.95 424 (68) 288 (29) L8,088,008+ itunout - test terminated

695-6 IVlr-291 SmOOth 8.95 379 (55) 166 (27) LS,888,0OG+ Runout - test te_mineted

582-4 IVlr-291 58ooth 8.55 L375(55) 166 (27) L860886888 - test terminated 

]

Heat treatment (:onststed of 1286"C (2358*r) fol: 2 hours, plus 1296"C (2365er) for 4 hours

vith repL.d Argon Quench, plus 962"C (1888"F) for 5 hours, plus 871"C (1600*,) for 28 hours.

k Ratio = alternating stress amplitude

IJeZ_ |tKeSS ._r

Frequency - 68 Hertz

approximately perpendicular to the applied stress in the (001)

plane and then switched to Sta_e I and propagated on a crystallographic plane. This trend is indicated in Figure 62 for an SC

HASAIR 100 specimen that failed at 6,477,000 cycles. Close examination of the casting pore associated with initiation also indicated that a small amount of crystallographic cracking was associated with initiation, but this condition was not typical.

In both HCF and LCF tests, cracks commonly propagated in

Stage II and then switched to Stage I. This trend contrasted with

fatigue test results for equiaxed and columnar grain DS superalloys

where the reverse trend was normally observed.

128

i
ORIGINAL FAdE ;_

OF POOR QUALITY

552190} _ =

0_=,=_=,,, ='m'-'-. ==..=..., qi,,,,

=,< 276(40i --'--"--' SCNASAIRIO0 "" "" .......

'--="_ SCALLOY3

..... DSMAR-M247

%

1361201- A=MAXIMUMSTRESSAMPLITUDE -: MEANSTRESS

I I I I I l I I 1 I I l 1

0 , , I 3 , ,,,f , , , , I I_'_I = I I I I ill '

• I04 107 J

": CYCLESTOFAILURE ',

Figure 61. Effect of A Ratio on HCF Capability at 871°C ,_

(1600°F)"

Leverant and Gell, 15'16 in studies of fatigue crack propagation in DS and SC MAR-M 200, determined that Stag•e I cracks propagated in a stable fatigue mode on {iii} planes and tended to propagate preferentially, but not exclusively, in the <ii0> directions.

4_

These investigations also showed that crack initiation and proFaga- _

tion on the {lll} slip planes of SC MAR-M 200 were favored by lower

temperatures and higher frequencies, while crack propagation in

Stage If, perpendicular to the applied stress was favored by ,"

higher temperatures and lower frequencies. 17 .':

P_

The extent of Stage I and Stage II cracking in SC NASAIR i00 _

and SC Alloy 3 is partially inconsistent with those conclusions for ;";

SC MAR-M 200. Stage II cracking can be a significant portion of :_

the fracture surface in both low temperature LCF tests and inter- _I

mediate temperature HCF tests, i_

| ,

129 _
ORIGINAL PAGE iS

OF POOR QUALITY

871°C (1600°F)/552 MPa (80 KSI)/A = 0.95/60 Hz

)l

;_.._ ---r'_ _'_,: _ ", , ._

i

hi.

200X 10X

INITIATION SITE: CASTING POROSITY INITIAL PROPAGATION: STAGE II

FINAL PROPAGATION: STAGE I --

LIFE = 6,477,000 CYCLES

Figure 62. HCF Fracture of SC NASAIR i00.

Reasons for the occurrence of Stage II cracking p:ior to

Stage I cracking are also unclear at this time. Purushothaman and

Tien 18 examined slow fatigue crack progagation in [001] oriented SC

MAR-M 200 and verified that Stage II cracks were propagating by

slip on multiple (iii) planes. On the other hand, _ddananda and

Shahinian 19 have recently presented evidence that fatigue cracks

, can propagate in Udimet 700 in the (901) plane in a cyclic pseudocleavage mode.

Notched Tests

Notched HCF tests were performed at 871°C (1600°F) with an A

ratio of 0.95 (pulsating tension) and a frequency of 60 Hertz (3600

!

130

®
ORIGIr4/%LP;_G_ _

OF POOR QUALITY

cpm). As indicated in Figure 60, the 6.35 mm (0.25 inch) gauge

diameter was enveloped with a notch with a stress concentration

factor, Kt, equal to three.

Results of these tests are provided in Table 56 and indicate

that both SC NASAIR 100 and SC Alloy 3 have equivalent 107 cycle

• run-out lives at a maximum stress of 345 MPa (50 ksi) or an alternating stress amplitude of 168 MPa (24.3 ksi).

Temperature

HCF tests were conducted at both 871 ° and 982°C (1600 ° and

1800°F) to establish the effect of temperature on HCF stress capability. Data obtained for SC NASAIR 100 and SC Alloy 3 at 982°C

(1800°F) and an A ratio of infinity are provided in Table 57. A

7[

TABLE 56. NOTCHED (K t " 3.0) HIGN-CYCLE-FATIGUE LIFE OF UNCOATED 6.35 MM (0.25 INCH)

+ DIAMETER SC ALLOY SPECXMENS.

s_cL.e,Mster * Ta,p,ratur,| i.*tsr,,tln9to

N_ber..... [ _ _IRstlo "C _'r) I Xss*,,.J A.p*It_s rs*Zu,e Sen,[ks _t

544-3 Vf-220 SC liAakZR 100" 0.95 071 (L600) 483 (70) 234 (34) 20,000

570-37 VP-235 0.95 448 (6S) 221 (32) 116,000

544-5 VF-220 0.95 440 (65) 222 (32) 255,000

522-21 Vr-220 0.95 414 (60) 200 (29) 57,000 i

$69-26 VF-235 0.95 i 414 (00) 200 (29) 1,017,000

569-29 VF-235 J 0.95 i 379 (SS) 100 (27) --- HachLne ss2funotLon

i

56$70-4 0-18L VF--2235 0.95 . 379 ]79 (SS)55) L1068 (27) 03394,000 ,000 PocoaLty •

$67-9 Vr-235 0.95 379 (SS) 106 (27) 1,176,000

567-12 VF-235 0.95 345 (50) L6S (24) L0,000,000+ Runout - teat termLnated

560-15 VF-235 p 0.95 J45 (50) 165 (24) L0,000,000. Runout - teat terstnetsd

68S-S VF-29L SC &2lay 3'* 0.95 440 (6S) 221 (32) 204,000

705-8 VF-29L 0.$S 414 (60) 200 (29) $1S,000

'_ 689-1 Vr-2gl 0.95 379 (55) 186 (27) 500,000

i 582-6 vr-292 0.95 379 (as) L06 (27) 1,360_000 _t

i 709-10 Vr-292 0.95 345 (50) L65 (241 L0,000,000 - test tersLnated , 

502-8 vr-291 l 0.95 ]45 (50) X65 (24) L0,000,000 - teat tscsLnated " 

!

• Best treatment consisted st L316"C (2400"f) for 2 ho_,rs, plus 13240C (24150r) for 2 hours

vLth r_pLd Argon Ouenoh, plus J820¢ (10000r) |or S hauls, plus 071*_ (2600*r} for 20 _)ute.

• **Best traitieht _oneLsted _r L260"C (2350*r) for 2 hours, plus 1296"C (2305"r) toc 4 hours

wLth [apLd Argon Quench, plul 9820C (1000or) tar 5 hours, plus 0710¢ (1000"F) for 20 hours. :_

" a_net_nu stress @!_Litude

A Ratio • • meen stress _i

rrsquen+y • 60 Hertz

131
ORIGINAL PAGE |g

OF POOR QUALITY

TABLE 57. HIGH-CYCLE-FATIGUE RESULTS FOR UNCOATED SC ALLOY 6.35 )04 (0.25 INCH)

DIAMETER SPECIMENS TESTED AT 982"C (1800*F).

s Bur f_e COnt lg- A Tealpereture ; [ AIt •C"_"_4t l._ to

55_-5 120 _ HABAI_ 100" Um:o4ted _lmoth . 012 (IJ0ll 34S (SOl 14S (SOl 967,000

SIT-10 _Vff-lJ5 Unooeted 8ROOth . ]L0 (4S) ]L0 (4S| 4,0]0,000

I

Sil-14 VP-2]S Un_oe_d SaooCh . 210:40) 271 (40) 4tL,000

569-30 VP-2]S UrmOated Slmoth w 276 (40) 276 (40) 7,?03w000

S70-3g VP-235 UncC_ted 811ooth e 242 (35) 241 (35) LO,O00,O00+ tunout -

test ¢e¢minated

549-9 VF-220 UncOated 8oooth . I 107 (J0) ]07 (30) 10,0000000+ lunout -

I feet te[aAna¢ed

6K6-1 Vr-291 SCAlloy3"" UncOated 811K_th . i 2?4 (40) 270 (40) 470,000

709-2 V]P-191 [ _ Uncured Saooth - 241 (35_ 241 (35) 3,343,000

595-5 Vr-291 ; UncOated Sil_tb - i 207 (30) 20? (30) 10,000.000+ hur_ut -

I . , , ; ,J . . . teet terminated

• Heat treetnent ¢o¢ SC :¢_AIR 100 ¢one_sted ot 1316"C (2400*F) |or

2 hou¢e plus 1324mC (]4_Se?) toy 2 hou¢l vAth [epld Aton (]ue_h,

pLul 91J'C (ll00"f) _oc S houri, Blue l?l*C (1400"Y) Eo¢ 30 houce.

seE/it ¢[eetMht _o¢ _ AlLoy $ ¢oniiit_ el 1288*C (2350"r) toy 2

hou[e plUS 1294"C (23iS*F) f_ 4 h¢_l[e with RapLd A¢gon Quench, plu_

902"C (1000"r) toy 5 hour0, plus 071"C (1400"r) lot 20 heuce.

A RetLo • Alterfllt_nqManS¢¢eseStleeeAmP!_cude , Y¢eqoency • 0.33 _e[t$

compacison with similar data obtained at 871°C (1600°C) can be made

usin£ Tables 54, 55, and 57. This comparison indicates that the _

nigher test temperature resulted in approximately a 3C-percent ""

reduction in HCF run-out stress capability for both alloys at the

higher Lemperature.

Coatings

Protective coatings are essential to the long-time surface

stability of superalloys in a turbine environment. However, under

some conditions of cyclic strain amplitude and coating thickness,

cracks initiating in the coating can propagate into the superalloy.

Thin coatings generally show the least tendency to penetrate into

the superalloy. 14'20

• To quantify any coating effects on the HCF capabilities of the

SC alloys, smooth 6.35 mm (0.25 inch) gauge diameter HCF specimens

were diffusion aluminide (RT-21) coated to a thickness of 25 to

50 _m (0.001 to 0.002 inch) by Chromalloy American Corporation in

Or_ngeburg, New York.

132
After coating, HCF tests were conducted at 871 ° and 982°C

(1600 and 1800°F) and A ratios of 0.95 and infinity. Test data for

SC NASAIR 100 and SC Alloy 3 are provided in Tables 58 and 59,

respectively. In addition, a comparison with coated and uncoated

data for SC NASAIR 100, SC Alloy 3, and DS MAR-M 247 is provided in

Figure 63 for an A ratio equal to infinity.

Inspection of these data indicates that the coating effects on

the 107 cycle life capability of the alloy were not significant at

the coating thickness evaluated.

Examination of the 107 cycle run-out specimens, however, indicated in all instances that the coating had initiated numerous

cracks in both alloys. Photomicrographs of diffusion aluminide

coated SC NASAIR 100 after 107 cycles of testing at 871 ° and 982°C "_

._ (1600 ° and 1800°F) are provided in Figures 64 and 65 to illustrate

$: the extent of coating cracking at runout. At both temperatures,

coating-initiated cracks were observed to propagate a short dis-

,j - -_

tance into the alloy. HCF conditions that would not initiate coat- ! ..

ing cracks were not determined.

Long-Time Exposure

!

A limited number of HCF tests were conducted at 982°C (1800°F)

%

with SC NASAIR 100 and SC Alloy 3 specimens that were machined from

fully processed bars that had a subsequent 982°C (1800°F)/1000-hour

furnace exposure. As noted previously, this exposure coarsens the

gamma prime phase in both alloys and results in precipitation of

: the mu phase in NASAIR 100.

|

¢

•" Results of tests for the SC alloys in the exposed condition t are provided in Table 60. Figure 66 is a comparison between •

,_ and data. of these data indi- exposed typical unexposed Inspection

' cates that the furnace exposure had no effect on the HCF strength %

of either alloy.

_- 133
ORIGINALPAGEiS

OF POORQUALITY

TABLE 58. HIGH-CYCLE-FATIGUE CAPABILITY OF DIFFUSION ALUMINIDE (RT-21) COATED 6.35 MM

(0.25 INCH) DIAMETER SC NASAIR I00 SPECIMENS.

..... 8tres"8-MPa (Xsi) Cycles

.Num p.bie.er ° ,Hea a.., t Configuration _acio _aximumJ Amplitude Failtureo Remarks

.. _ . .

570-43 VF-235 Smooth ).95 87L (1600) S89 (10031 338 (49) 41,000

589-36 vr-233 SJKx)th ).95 689 (100) 338 (49) 301,000

544-14 vr-220 Smooth ).85 621 (90) 303 (46) 546,000

567-3 VF-235 SrJoth 3.95 621 (003 303 (44) 1,180,000

544-25 VF-220 Smooth ).95 586 (85) 283 (42) 2,894,000

552-13 Vr-220 6mooth _.85 552 (80) 269 (38) 6t477,000

553-23 VF-220 Smooth _.95 552 (80) 289 (39) 8,455e000

552-24 VF-220 Smooth 10.05 483 (70) 234 (34) 2, OO,OOO

568-21 VF-235 Smooth _0.95 483 (70) 234 (34) 4,070,000

568-23 vp-23_ Smooth 10.95 483 (70) 234 (34) 10,000,000+ Rumour - teat terminated

i

553-8 vr-220 8mc_th I 0.95 448 (65) 221 (32) '10,000,000+ Rumour - test terminated

1

569-32 VF-235 J Smooch 0.95 448 (65) _21 (32) LOsOO0,000+ Rumour - teat terminated

544-11 vr-220 Smooth - 552 (80) 552 (80) 36,000

570-45 vr-235 8_ooth - 517 (75) 517 (75) 36,000

544-22 VF-220 Smooth - 483 (?0) 483 (70) 48,000 J

56g-31 VF-235 I 8neath . 483 (70) 463 (703 118,000

568-20 Vff-23_ 8_th _ 414 (60) 414 (60) 350,000

552-11 VF-220 Smooth _ 414 (60) 416 (60) 484,000

567-6 vr-235 SnoGth - 379 (55) 379 (55) 3,512,000

553-28 VF-320 8aooth . 379 (55) 370 (55) 5,428,000

570-42 Vr-235 Smooth . 365 (533 365 (533 482,000 Porosity

552-18 V_-220 8imoth - 345 (50) 345 (50) 15,000 Extensive canting

porosity

553-22 Vr-220 Snooth _ 345 (503 345 (S0) L0,0OO,00O - test terminated 

567-1 VF-235 $moot_ - 345 (50) 345 (50) 2O,0O0,00O+ Rumour - teat terminated

544-16 VF-_O Smooth - 982 (L800) 345 (503 345 (50) 27,000

553-27 VP-220 Smooth - 310 (45) 310 (45) 7,660,000

'569-33 Vr-335 Snoot% _ 276 (40) 276 (40) 7,563,000

! 567-5 VF-235 $nK)Oth " 276 (40) 276 (40) 8,L66,000

552-17 VF-220 Smooth - 276 (40) 276 .(40) 10,000,000+ Rumour - teat terminated

568-22 VP-2_5 Smooth " 24_ (35) 241 (35) 10,000,000+ Rumour - teat terminated

float treatment consisted at 1316*C (2400"r) _or 2 hours, plus L3240C (2415er) tar 2 hours

vith rapid A_gon Quench, plus RT-_L c(_ttad, plus 8720C (1600*r) _or 20 hours.

_lternet|n_ ,tress 8mplitqda

A Ratio " ' mean Italia

rKequency • 60 Hertz

134
ORIGINAL PAG_ ;$

OF POOR QUALITY

TABLE 59. HIGH-CYCLE-FATIGUE CAPABILITY OF DIFFUSION-ALUMINID}: ,RT-21) COATED

6.35 MM (0.25 INCH} DIAHETER SC ALLOY 3 Si :IMENS.

S

N

pe

um

cl_

b

m

er

J

I

Nes

He

t

at

er J

I Con_AguratLon Jdn

k

tlo iTemeC_

per

(elP)

ature

NalLm**un

i

_

_i"*-

A

to

iq

rn

p

etln

1Ltud

g

e 1_Fa+°+toLluce *''_-__-'_'K8 .... r

695-4 VF-291 Smooth 0.g5 I 071 (1600) 414 (10) I 200 (29_ I 34,000 Unacceptable _[oaLt¥

709-6 VIP-291 i S_th 0.g5 I 414 (60) I 200 (29)

582-5 VIP-2gl I Szk_th 0.95 I 379 (55) I 186 (27) I 140,,000 021,,000 600+ Surtout - test ternAneted

582-:1 VF-291 ; S_th - 310 (45) m 310 (45) 1,387,000

.: 695-1 VF-291 Sw.._th m 310(45I) 310

709-3 vF-291 Smooth - 276 (40) I 276 (45) 3,492,000 • (40) 10,000,000 Sunout -telc termAnnted

709-4 Vg-291 I Smooth "q 902 (1800) 310 (45) 1 310 (45) 96,0u0

695-3 VF-291 , Smooth " ' I 276 (40) I 276 (40) 2,175,000

582-3 VF-291 j S_th w v _ 241 (35) I 241 (35) 10,O_,u_O', itunouc - tilt terminated

Heat treatment consisted ot i288*C (2350*W) for 2 houri plus 1296*C (2365"F) for 4 hours _tth Argon Ouench, i)lul ST-21

coating, plus 871eC (1600°F) (or 20 hours.

AlternetJnq_treesAm_lltude

Ratio • Nean Stress

Frequency • 6Q Herr7

•" A = oo/60HERTZ

483170)- 871oC(1600° F] ,mira, 982°C(1800°F) _

--414160] 414(flOl345(50J- _'_

:S _ A ---.. -.--,., _=Z 3<31501-

r,,,5 2761401 ......... [] 278(4Ol- E)

,1, -'_'_-- _ _ _1 "

_ 207i301- _ 2071301-

a¢ t38120l 4- <3E 1381201-

6911_1j 001101-

' I : tl 0 t ' t I _--; -I _ ''

,* 106 107 108 lrj7

, CYCLESTOFAILURE CYCLESTOFAILURE ',,

•( UNCOATED RT-21COATED ',

SC NASAIRI00 C)

•i--, SC ALLOY3 ,/_

.... OS MAR.M247 _ _+.

_ Figure 63. Diffusion-Alumlnlde (R','-211Coating Effects

i on HCF Properties. '

q

135
oRiGINAL PAGE iS

OF PoOR QUALITY

t

l

• °

t_

. RT21COATING-- ml__"']'_"l 1"_'"_"_ _'_'_-"_'_-_'='_."_:_',_'_"";_%_"

I

i•

°

i'=

-'" SCNASAIR100

:" 80X

PSOSZS.2

_.."

RT21 COATING :

,o

I

,- SCNASAIR100

320X

P80825-1

Figure 64. HCF Cracks in RT-21 Diffusion-Aluminide-Coated

SC NASAIR 100 Specimens After 107 Cycles at

;, 871°C (1600°F) and a Maximum Stress of 345 MPa

" (50 ksl). (A ratio equal to _).

; 136
ORIGIr_ALPJ\C[;_,5

OF POOR QUALITY

o

%, , . .....

-- _ RT21 COATING

t

SCNASAIR100

80X

"_ P_25-4

"

RT21COATING .:4

1

SCNASAIR100 :.'

e

"1

320X "

P80825-3 ;.

| ' I

Figure 65. HCF Cracks in RT-21 Diffusion-Aluminide-Coated ',J.

, SC NASAIR i00 Specimens After 107 Cycles at [_

982°C (1800°F) and a Maximum Stress of 241 MPa ;.

(35 ksl), A = oo. _,

't

137

.i

4
ORIGINAL PAGE 19"

OF POOR QUALITY

TABLE 60. HIGH-CYCLE-FATIGUE RESULTS FOR AXIAL-AXIAL TEST OF UNCOATED 6.35 MM

(0.25 INCH) DIAMETER SPECIMENS MACHINED FROM EXPOSED t SC NASAIR I00 AND

SC ALLOY 3 SEPARATELY CAST TEST BARS.

t I.... I Itrusjlah _ " Cycles

Specimen Nester Con_lg- I A Ten.stature Alto[noting t.o

_-_,_,-•heSr est.L._._11o__ ur==lo.Imac. so] *c ('r) N==*,,u,, A=lpllt.d* WaLIuro a_s=ks

653-1 VF-236 I_ =ASAZR 100 S=ooth I " 902 (1000) 310 (45) 310 145) _t,20S,000

653-2 VlF-236 I 843 IIASAI'R 100 Smooth I 0" 276 140) 276 (40) 6,9]L2,000 I

653-3 Vr-236 ISC NASAIR 100 Gmooth J "0 241 (35) 241 (353 .0,00000004" Runou LefmJnmt-tssf,. tsd

60G-t VF-2gL I SC:Alloy 3 Smooth I " 276 (40) 276 (40) 3,124,000

66&-2 VF-291 I SC Alloy 3 Smooth I '_ 241 (35) 24.1. (35) 1,614,000

666-3 V¥-291 I .SO Alloy 3 Smooth I ." 1207 (30) 207 1303 10,000,0004" Runout:-tsst te,ninal:sd

i

[ L

*Pc|or to exposure at g82"C (1000*r) _o_ 1000 hours=

SC HAgAIK 10Q heel treatment consisted of= 13160C (2400*r) Co( 2 hours# plus 1324"C (2415"r) for 2 hours

vith cspid &rqon Quench, plus 982"C (1800"r) lot 5 hou:s, plus 071"C (1600"r) Eor 20 Lou_s. ._r_

SC A11o¥ 3 heat t[eatment consisted of 1288"C (2350"F) foc 2 hours& plus 1296"C 12365"r) for 4 hou:s

wLth rsp_d Argon Ousnch, plus g02"C (1800"F) to[ 5 hou_8, plus 071 C (1600 r) £oc 20 hours.

h Ra_lo " alternating stress amplitude Frequency • 60 Hertz

' ' mean stiesJ

6.3 Environmental Tests

Successful operation of SC turbine blades is dependent on

maintaining airfoil surface integrity under conditions of high temperature oxidation and intermediate temperature hot corrosion (salt

film accelerated oxidation). Consequently, burner rig testing was

conducted to assess the environmental resistance of uncoated and

RT-21 diffusion aluminide coated SC NASAIR 100 and SC Alloy 3 specimens (Figure 67). Coatings were applied to an approximate thickness of 50 _m, (2 mils) by Chromalloy A_=erican Corporation in

Orangeburg, New York.

Hot-corroslon testing was performed on duplicate specimens of

uncoated and RT-21 coated SC NASAIR 100 and SC Alloy 3 at 927°C

(17000F) for 307 hours. Burner rig test parameters included the

following:

o Cycle - 60 minutes hot + 3 minutes forced air cool

o Fuel - Jet A

138

®i
C

ORIGIKAL _"-'-'*"

OF POOR GUALFFY

982°C(1800°F)/A= /60 HERTZ

' 483170)

414160)-

.. __.345150_} SCNASAIR100

,.,. SCALLO3Y _ _.. _. /k , ,

_oo2071301- IFULLPROCESSED) Y _" _ _ _:

X

,mC

138(20_--

:._'" 0 SCNASAIR100} EXPOSEDAT982°C(1800°F) :. 69110J- /k SCALLO3Y FORlO00 HOURS

"" _

o I I t I I I tt _

- 106 I07

? CYCLESTOFAILURE

Figure 66 Effect of Long-Term Exposure on HCF Strength

76.2mm £.- 12.5mrn "_

' Jl= (3.000IN.) _ __i IN.) DIAMETER

" 50.8mm _ "' V

' 22.2mm 12.0 IN.)

I (C.875 IN.) ,_

__ 15° ,,

4.7mm ;1

r_

10.187IN.) .4 ,

,_, Figure 67. Burner Rig Specimen t

.t 139
ORIGINAL PAGE _S

OF POOR QUALITY.

o 5 ppm synthetic sea salt (ASTM DI141-52)

o Specimen rotation - 1500 rpm

A description of the burner rig is provided in Reference i.

The appearance of the RT-21 coated SC alloy specimens after

hot-corrosion testing at 927°C (1700°F) for 203 and 307 hours is

shown in Figure 68. The RT-21 coated SC Alloy 3 specimens exhibit

some oxide scale spallation without base-metal corrosion. The

RT-21 coated SC NASAIR 100 specimens exhibited coating failure and

base-metal corrosion.

.

0, ; I

"J ,

+o

i . • ++_."

+ t P74143-2

e ,'

c

SC ALLOY 3 SC NASAIR 100 SC ALLOY 3 SC NASAIR 100

263 HOURS 307 HOURS

Figure 68. Appearance of RT-21 Diffusion-Alumlnide-Coated

Specimens after Indicated Hours of HotCorrosion Testing at 927°C (1700°F) with

5 ppm Sea Salt. (Magnification: 2X)

140

1 ®
ORIGINAL PAGE IS

OF POOR QUALITY

Weight change data for hot-corrosion tested uncoated and RT-21

coated SC NASAIR i00 and SC Alloy 3 specimens are provided in Table

61. For comparison, the DS Mar-M 247 (uncoated and RT-21 coated)

927°C (1700°F) hot-corrosion test results generated in Project 1 of

the MATE Program are also included in Table 61. Based on weight

change and visual examination data, the test results indicated the

following:

• o Uncoated SC NASAIR i00 and SC Alloy 3 exhibited corrosion

similar to uncoated DS Mar-M 247.

o RT-21 coated SC NASAIR 100 exhibited less corrosion

resistance than RT-21 coated SC Alloy 3 or RT-21 coated

DS Mar-M 247. Coating failure of RT-21 coated SC NASAIR "_:

_ i00 specimens occurred after 203 hours.

'l

TABLE 61. 927°C (1700°F) HOT-CORROSION TEST RESULTS.

t

t Weight change* (zn grams) at _ndzcated test txme (hours) I Remarks _:

.... 25 50 76 I01 I15]" I20J 1256 L307 1Done °

Alloy - _ . . :

NASA[R 100 -0.07 -0.28 -0.87 -1.47 -2.00 -3.41 -4.57 -6.09 Gross cor_oslon _'

(Uncoated)

NASAIR i00 +0.01 +0,01 +0,01 0 -0.01 -0,02 -0.04 -0.25 Coating failure

(RT-21 Coated) after 203 hou:s

Alloy 3 -0.02 -0.07 -0,45 -1.04 -1,61 -2,73 -4.01 -6.75 Gross co_oslon

(Uncoated)

,\llo_ 3 +0.01 +0.01 +0.01 i+0.01 +0.01 +0.01 No coatlng 

(RT-21 Coated) deg -_adat _on

t _- [. i

' i

_, 20 40 80 i00 150 210 260 310 :

PS Ma_-M 247 -0.06 -0.13 -0.56 -1,09 -2.06 -5.14 -6.63 -9.83 Gross corrosion ._

Uncoated )

i DS Mat-M 247 0 0 0 0 0 +0,01 0 +0.01 Sl_ght coating (RT-21 Coated) degradation _'|_

*Wezgnt change is an average of 2 l_est specimens

I

., 141

' i
o RT-21 coa_ed SC Alloy 3 exhibited corrosion resistance as

good as, or slightly better than, RT-21 coated DS Mar-M

247. No visual evidence of coating failure was observed

on Alloy 3 specimens after 307 hours, whereas the DS

Mar-M 247 specimens exhibited slight coating degradation

after 310 hours of testing.

Burner rig oxidation testing was conducted on duplicate specimens of uncoated and RT-21 coated SC NASAIR 100 and SC Alloy 3 at

I093°C (2000°F) for 509 hours with the following test parameters:

o Cycle - 60 minutes hot + 3 minutes forced air c9oi

o Fuel - Jet A

o Specimen rotation - 1500 rpm.

: The post-test visual appearance of the RT-21 coated SC Alloy

specimens is shown in Figure 69 and weight change data for both

uncoated and coated specimens are provided in Table 62. "_

_ Examination of these data indicates that uncoated SC NASAIR

I00 appeared to have slightly better oxidation resistance when compared to uncoated SC Alloy 3. Weight losses for these SC alloys in

this I093°C (2000°F) test were significantly greater (by a factor

of 20 to 30) than those observed for uncoated DS Mar-M 247 specimens evaluated for 510 hours at I038°C (1900°F) in Garrett's NASAl

MATE Project i.

RT-21 coated SC NASAIR i00 and SC Alloy 3 appeared co exhibit

comparable oxidation resistance based on weight change data.

Visual inspection of the specimens indicated that the RT-21 coated

SC Alloy 3 specimens appeared to be degrading by a pitting process,

while the RT-21 coated SC NASAIR I00 was degrading by oxide scale

spalling (Figure 69).

- •

' 142

_ ®
ORIGINAL FAG_ I_

OF POOR QUALITY

I

•" P7414_7

L

SC ALLOY 3 SC NASAIR 100 SC ALLOY 3 SC NASAIR 100

Figure 69. Appearance of Duplicate RT-21 Coated SC

NASAIR 100 and SC Alloy 3 Specimens

after Oxidation Testing at I093°C (2000°F)

for 509 Hours. (Magnification: 2X) i

TABLE 62. 1093eC (2000*F) OXIDATION TEST RESULTS.

weight Change* (In graiI) ,;t indicated temt time (In hours)

.**o. -i so,I" I *°°I *'*1 _**[,'*_ I '_ j_o,

NASAIR 100 e -0.02 -0.02 -0.03 -0.04 -0.05 -0.05 '0.12 -0.16 -0.22

(Uncoated)

_ NASAIR 100 0 0 0 -0.01 -0.01 -0.02 -0.04 -0.06 -0.10 ,

(RT-21 Coated)

AllOy 3 -0.02 -0.04 -0.06 -0.07 -0.10 -0.14 -0.21 -0.26 -0.34

(Uncoated) ,_

_11oy 3 +0.01 +0.01 0 -0.01 -0.01 -0.03 -0.05 -0.06 -0.08 .

(RT-21 Coated) 0_

*Weight change le an ave:age o_ 2 telt |F4ol:enl

;

2
ORIGINAL. pAGE. tS

OF PoOR QUALt1"Y

Again, weight losses for RT-21 coated SC alloys were significantly greater (a f_ctor of 8 to i0) after 509 hours of testing at

I093°C (2000°F) relative to RT-21 tested for the same time at

1038°C (1900°F). This result is attributable to both the higher

test tem@erature and elimination of all or most of the hafnium from

the SC alloys.

• 6.4 Physical Properties

6.4.1 Thermal Expansion

Thermal expansion measurements were performed by Southern

Research Institute in B_rmlngham, Alabama. Test data were obtained

for triplicate specimens (Figure 70) of SC NASAIR i00 material

using the quartz dilatometric technique. Measurements were taken

k_w

_15.Ilmm

e----,e -- ('-_/l IN.)

/

7S.2mm (3 IN.) SPHERICAL RADIUS

t, ON BOTH ENOS

Figure 70. Thermal Expansion Specimen.

144

®
ORIGINAL FAGE IS

OF POOR QUALITY

every 56°C (100°F) from room temperature to I093°C (20J0°F). The

thermal expansion for SC NASAIR 100 is presented in Figure 71. As

expected, thermal expansion data for SC NASAIR 100 is virtually

identical to that for DS Mar-M 247.1

6.4.2 Thermal Conductivity

Thermal conductivity measurements were performed by Southern

Research Institute. Conductivity data were obtained for triplicate

specimens (Figure 72) of SC NASAIR i00 using the comparative rod

apparatus technique. The thermal conductivity measurements were

taken over the range of 93° to 1093eC (200° tO 2000°F) and are provided in Figure 73. Comparison of the SC NASAIR 100 conductivity

data to the DS Mar-M 247 conductivity data I revealed no significant

difference in thermal conductivity between the two alloys.

6.4.3 Elastic Modulus _ _r

i

Dynamic modulus of elasticity data for SC NASAIR 100 specimens

were measured by Southern Research Institute (Birmingham, Alabama).

_'o(25)

120)

z 115) ...... ._

X

.a (10) .....

in

_- 15) ....

00 _ 93 204 .. 316 427 538 049 760 671 962 1093 1204

(200) (400) (600) (800) (1000) (1200) (1400) (1600) (1800) (2000) (2200)

TEMPERATURE, °C (OF)

Figure 71. Thermal Expansion of SC NASAIR i00.

145

®
25.4 (1.,000) i

•o.o25 ;

3.18 _ 3.18

_9.125) " (0.125)

15°

15° , , _,

18 +0.025 (+0.001) .,

25.4 (1,000)

1.70 (0,067) DIA., (3 PLACES

NO. 51 DRILL)

=,

%.

p - _ - ! - m 25.4

< (1 000)

NOTES:

: ALL DIMENSIONS ARE IN

16 MM (INCHES)

+0,127 (+0.005)

; EXCEPT WHERE NOTED

Figure 72. Thermal Conductivity Specimen.

i

J 146

_m
ORIGINAL F "_".... + "*'

_t _,% _.,..3

OF POOR QUAI.iTY

_1._ (2_)

>_ 1.12 (2_)

pz

Zm

_ 0._ (100) _I/

o.=8m)

0

< 0 93 2_ 316 427 538 .9 760 871 982 1_3 12_

(200) (400) (6_) (8_) (1000) (1200) (1400) (16_) (1800) (2_0) (2200)

TEMPERATURE, °C (OF)

._ Figure 73. Thermal Conductivit_ of SC NASAIR 100.

Test specimens were machined from bar [001] and slab [iii] castings. The test specimen configuration for these tests was a +

rectangular plate 10 cm long, 1.3 cm wide and 0.13 cm thick (4 by

I

0.5 by 0.05 inches). Tests were conducted at room temperature and

at 93°C (200°F) temperature increments from 538°C (1000°F) to

982°C (1800°F). Each specimen was measured three times.

The modulus results for fully processed SC NASAIR I00 speci- _

mens oriented in the [001] + 2 degree and the [iii] + 5 degree

directions (determined by the X-ray Laue technique) are provided in

Table 63 and Figure 74. Examination of the results indicates the

i

following:

+i

o [001] specimen. The room temperature modulus measured on

I SC NASAIR i00 was I14.5 GPa (16.6 x 106 psi), which was _

I slightly lower than anticipated and compares with the

' published value of 125.6 GPa (18.2 x 106 psi) for pure

+ nickel In the [001] direction. '

i 147

!

r
ORIGINAL PAGE _,S

OF pOOR QUALIW

TABLE 63. DYNAMIC MODULUS OF ELASTICITY OF SC NASAIR i00.

6

MODULUS OF ELASTICITY, GPa (E X i0 PSI)

Specimen TEMPERATURE, =C (°F)

Orientation 24 538 649 766" 871 982

(75) (1000) (1200) (1400) (1600) (1800)

.. .,,, •

[001] + 2" 114.5 108.3 102.8 98.7 90.4 83.5

(16.6) (15.7) (14.9) (14.3) (13.1) (12.1)

[iii] + 5 ° 280.8 260.1 251.8 240.1 231.1 215.3

(40.7) (37.7) (36.5) (34.8) (33.5) (31.2)

b ......

. NOTES:

a. Values are an average of three readings.

b. Specimen configuration: i0 cm long, 1.3 cm wide, 0.13 cm

_ thick (4.0 by 0.5 x 0.050-inch).

o [ill] specimen. The room temperature modulus measured on

SC NASAIR 100 was 280.8 GPa (40.7 x 106 psi), which is

slightly lower than the published modulus value of

;" 293.9 GPa _42.6 x 106 psi) for pure nickel in the [Iii]

" direction.

o A comparison of the data indicates that the elastic

modulus of the [ill] specimen was about 2.5 times greater

_han that of the [001] specimen over the temperature

range examined.

148

.......... .. ®
ORICiNRL ?,ij;;' _,'

OF POOR QU._,L{T'(

-50

300-

_ - 40

= -30 =" 200

"-' 1106P_I] ..,J

.,%

-: _ -20 • o

.'=,,,--I 100

LU

-10 i:

0 | II l_ | | l | | | ]1 i| O

0 500 (932) 1000(1832)

- TEMPERATURE °C (.°F)

,, Figure 74. Elastic Modulus of SC dASAIR t00 in [001]

|, and Jill] Directions.

f: ,

149

®
I

?

' SECTION VII

7.0 BLADE DES IGN

{

L This task included the design activity required for the devel- L

_: opment of an uncooled, single-cryst_l (SC) high-pressure turbine

_ • blade for the TFE731 turbofan engine, using the alloys character-

[: ized in the previous section. Two blade designs were established

_ in this effort--the preliminary (initial) design and the final

" design.

< The preliminary design (PD) was established early in the pro- "_

_ " gram to provide a blade casting design suitable for use in the

_ development of the exothermic SC casting process and associated

i_ " material evaluations. This design was based on the final design

2 . of _he DS blade from MATE Project I. _

i

.[ Actual material properties and other data obtained from pre-

{ liminary design blades cast in SC NASAIR i00 and SC Alloy 3 were

= _ed in establishing the final blade design. The geometry of this

final design made it necessary to modify the turbine disk, nozzle,

and several other turbine components of the TFE731 engine to permit

;. effective integration of the blade into the engine assembly. The

; redesign of these turbine components was also completed in this

task.

7.2 Prel_minary Blade Desi@n

The preliminary blade design (PD) was established primarily

to provide a realistic blade configuration to be used in all the

° SC process and alloy development work scheduled in the first part '

of this project. This PD blade is the same configuration as the

_ final blade design established for the MATE DS blade in Project _,

: and, in fact, used the same wax tooling.

I
7.3 Final Blade Design

7.3.1 Mechanical Desi@n

Single crystal material is characterized by anisotropic material properties. The variance in the modulus of elasticity is

especially dramatic, varying from a low of 18 along one of the

principal cubic axes (001, 010, 100) to a high of 43 along the cube

diagonal (iii). This variation in properties will affect both the

stresses within the part and the vibratory response of the structure.

To accommodate these SC anisotropic material properties, ap- 5

propriate Garrett computer analysis programs were modified to include property variations in the input. These program changes al-

_ lowed both static and dynamic blade response (centrifugal and thermal stress plus vibratory modes and frequencies) to be evaluated.

%

• ._ $

" To study the effect of anisotropic properties analytically,

sample check analyses were run at various crystal orientations,

using the preliminary design blade configuration.

-"c

The finite-element model used for these studies, is shown in

Figure 75, and the frequencies calculated with this model are tabulated in Table 6 / . The calculated mode shapes for the blades at

• several natural frequencies and various crystallographic orienta-

; tions are shown in Figures 76, 77, and 78. The corresponding cal-

. culated stress distributions in the turbine blade at various crystallographic orientations are shown in Figures 79, 80, 81, and 82.

All stress and vibration studies for this preliminary design phase

were calculated at room temperature.

A comparison of these analytical results to the measured fre-

,_ quency response of various PD blades cast with a variety of crys- ;

_ tallographic orientations is shown in Figure 83. As can be seen,

a significant amount of scatter, attributed to blade-to-blade variations and limitations of GTEC holography capability at that I

152

: ®'i
ORIGINAL PAGE iS

OF POOR QUALITY

," e

_ Figure 75. Finite-Element Model (DS TFE731 HPT

i Blade Used for SC Material Studies).

TABLE 64. EFFECT OF CRYSTAL ORIENTATION ON RESONANT

FREQUENCIES (SC TFE731 BIGH-PKES3URE

." TURBINE BLADE).

Material

' Orientation (Degrees)

Prlmary (1) Secondary (2) Natural Frequ_ncles (Hz)

_' Is,_tropic 5504 9876 16463 21017 26471

;f* 0 0 5325 10436 15980 20864 26338

4

• 15 0 5175 10396 16227 20457 26144

_ 30 0 5173 9653 16321 19838 25366

, 0 -15 5319 10461 15944 20844 26379

0 -30 5317 10476 15930 20883 26350

(1) Angle between (001) material direction and blade Z-axis

* (stacking axis)

"_ (2) Angle between (100] m&taJial direction and blade X-axis

- _ (rotation axis)

k

,_ 153
ORIGINAL PAGE I|

OF POOR QUALITY

°,

L . }

i

.",r,

• MODEI FREO(HZ)= 5325 MODE2 FREO(HZ}= tO43B

- .j

1

i MOOE3 FREQ[_)= 15980 MOOE4 FRE_tHZ]=20864

Figure 76. Mode Shapes for Material and Global Axes

;., Aligned, RPM = 0, Room Temperature.

t is4 (_

t
I

_p

!

ORIGINALPAGF..19

OF POORQUALITY

p

/.

" -_ MODEI FREQ(HZ)=51"/3 HODE2 FREQfHZ]=9653

r ,b

2

br_

!

MOOE3 FREQ(HZ)-16321 MODE4 FRffQ{HZI1=9B3B ;q

• i

m

'_; Figure 77- Mode Shapes for 30-Degree Primacy Orientation,

',_ RPM _ 0, Room Te_,,perature,

-i ].55

:Ca ®'
ORIGINALp/_,GEI_

OF PoOR QUALITY

=, II

c

i'

t

i

'" MODEI FREQ{HZ)= 5317 MODE2 FREQ(MZ)=10476

• 0

i

_o

MODE3 FREQ(M'L}=15930 MODE4 FREO(HZ}=20883

g

: Figure 78. Mode Shapes fo_ -30-Degree Secondary Orientation,

RPM = 0, ROOm TemperatuEe.

156

I e,, _' ®
I

ORIGINAL PAGE 19

OF. POOR QUALITY

• PRESSURE SUCTION

• IDE DE

. :. Figure 79. Stress Distribution for Isotropic Material

:- RPM = 30,000, Room Temperatu_'e, • = Maximum Stress, • = Minimum Stress (ksi).

._

"- "4

4

J

PRESSURE SUCTION

SIDE SIDE

P

t '

Figure 80. Stress Distribution for Material and Global Axes ,

" _ Aligned, RPM = 30,000, Room Temperatqre, •= Maxi- I

' I

; I mum Stress, &- Minimum Stress (ksi). i

157
(JRIGINAL PAGE iS

OF POOR QUALITY

s,., SiDE c-,-,o,,,

:. Figure 81. Stress Distribution for 30-Degree Primary

. Orientation, RPM = 30,000, Poom Tempez'ature,

• Maximum Stress, • Minimum Stress (ksi) .

• j

]

PRESSURE SUCTION

S,DE , f_" _ sloe

',

. Figure 82. Stress Distribution for 30-Degree Secondary

_ Orientation, RPM = 30,000, Room _e_aperature,

' • Maximum Stress, • Minimum Str0- _ (ksi)

Ir

' 158
ORIGINALP_,C_ _;_

OF,POOR QUALITY

3O

O O

c_ d) '_,b TWO*

: o o o • o SECOND*

• • • o 'k • COMPLEX

20 _ -El- ,6 .... -- o_ SECOND

" o _J o TORSIONAL

. o •

-j,

,,., FIRST

cEL_i • 1" • 1" tt" r- COMPLEX

"OF I0 ,,, ,_ A . ,_ z_ ._. _ __ _ _ FIRST _'

A "_ "_ A _ A TORSIONAL

.-

" - k i• •_ FIRST

,"._ --"-m mw"1_ I1 -in -- iF .... 1-FLEX

: 0

0° 5° I 0° 15° 20° 25° 30° 35°

CRYSTALORIENTATION

LEGEND

• CALCULATEDFREQUENCIES, PRIMARYORIENTATION

• • CALCULATEDFREQUENCIES. SECONDARYORIENTATION

• FIRSTFLEX.MEASUREDFREQUENCY

, _ FIRSt TORSIONAL. MEASURECFREQUENCY

• FIRSTCOMPLEX. MEASUREDFREQUENCY

f o SECONDTORSIONAL, MEASUREDFREQUENCY

_' • SECONDCOMPLEX. MEASUREDFREQUENCY

_' o TWOSTRIPE.MEASUREDFREQUENCY " |

.:',r *EXCESSIVESCATTERPROHIBITSMAKINGANACCURATEESTIMATE

¢ OFA BESTFIT CURVE

I:

Figur e 83. Comparison of Holography Data and Analytical

_ Predictions of the Crystal Orientation • Effect on Vibrational Response.

"_ 159

I

4

A
ORIGINAL PAGE i_

OF POOR QUALITY

time, exists in the measured holography data. This occurs in the

higher order modes, such as the second complex and two stripe.

Even with this scatter, there appears to be no significant relationship between crystal orientation and blade frequency for the

range of orientations examined. Details of the measuled frequency

response of these blades are presented in the Table in Appendix A.

It should be noted that these results may be influenced by blade

size and configuration; and, therefore, this response may not be

typical of all single crystal blades.

Stress and vibration analyses of the SC turbine blades at

actual engine conditions were initiated using the model shown in

Figure 84. The blade configuration for the initial analysis is the i

same as the PD blade. Using the finite-element program previously d

modified to include anisotropic SC material properties, the radial 'i

_ and equivalent blade stresses were calculated as shown in Figures 85 and 86. (Equivalent stresses are the conventional method

: of relating triaxial component stresses to uniaxial material

strength.) i

J

Figure 84. Single-Crystal Blade Model.

160
ORIGINAL PAC_ tS

OF POOR QUALITY

_ PRESSURE _SUC',ON

!.!! SiDE _ E

_. _ 0 I

Figure 85. Calculated aadial Stresses for the Single-Crystal

. Blade, RPM : Maximum Takeoff, •- Maximum, _

-L & = Minimum (ksJ.) .

%.

%

• _4t

d _

_ T

;" PRESSURE _ SUCTION

• SIDE SIDE

2,

Figure 86. Calculated Equivalent Stresses for the SingleCrystal Blade, RPM = Maximum Takeoff

• = Maximum, • = Minimum (ksi)

; 161
7.3.1.1 Blade Life Analysis

Based on the results of this analysis, it was determined that

the PD blade configuration was potentially acceptable for the final

design. To verify this assessment, a detailed design blade life

analysis was initiated. This life was based on the planned MATE

200-hour test mission, which includes the four 50-hour cycles shown

in Figures 87, 88, 89, and 90. The normalized engine speeds _nd

scaled temperatures fo- each operating point are shown in the Table

below.

Cor_dition N 2 T 4

Takeoff NTO T 4 Mar

Max Continuous 0.990 NTO 0.975 T 4 Max

1 0.987 NTO 0.975 T 4 Max

2 0.985 NTO 0.975 T 4 Max

3 0.985 NTO 0.975 T 4 Max

4 0.985 NTO 0.975 T 4 Max

5 0.985 NTO 0.975 T 4 Max

6 0.985 NTO 0.975 T 4 Max

7 0.985 NTO 0.975 T 4 Max

8 0.985 NTO 0.975 T 4 Max

9 0.985 NTO 0.975 T 4 Max

i0 0.985 NTO 0.975 T 4 Max

ii 0.985 NTO 0.975 T 4 Max

12 0.985 NTO 0.975 T. Max

13 0.985 NTO 0.975 T 4 Max

Idle 0.608 NTO 0.299 T 4 Max

162

.... i

_o
ORIGINAL PAGE iS

OF POOR QUALITY

" 13 INCREMENTAL THRUST SETTINGS

MAX CONTINUOUS --m

1 1

2 2

3 3 3

4 HOURS 4

MINIMUM

e 8

Z 7I

I.-

I- 9 | 9 : IJJ

m _ 1.5 HOURS MINIMUM Io

-. ,-. 11 TYPICAL 11

LU 12 12

_) 13 IDLE 13

2.5 HOURS 2.5 HOURS _l

MINIMUM MINIMUM

D .'" ,till

I I I I

o-

_. . 0 19 20 30 40 % "*

"-" TIME - HOURS

.. Figure 87. High-Cycle-Fatigue Cycle A. _'

t

"" MAX CONTINUOUS

i i

"" 148 HOURS) '_

=

,,

(3

' z_

k.-

,1" P"

UJ

UJ

% 2 "

1. !

I

IDLE

',

• I I I :

o 1o 2o 3o 4o 50 "_

j TIME-HOURS '_

t

. Figure 88. Stress-Rupture Cycle _.

&

: 163

•1

A ®
ORIGINALPAGE [_

OF POORQUALITY

TOTAL 120 CYCLES

---_ _o- 10 MINUTES

!

MINIMUM

•, MAX CONTINUOUS

'I

|

t

-, _--- m 10 MINUTES

:" M|NIMUM a

IDLE _'_

2S MINUTES mm /CYCLE m_

,t A

)" I W

,- _ _ 20 _ ;o 50

" "rIME - HOURS

j,.

' Figure 89. Low-Cycle-Fatigue Cycle C. ...

i

; lOOCYCLES=> I'OTAL50 HRS

t;,o "l _ 11MINI

,, '-'IMAx CO_T(3.5MiNI

' _,RUISE(20.5MINI

'.. . • ................

a¢

•_ I,-,,

t-,-

, LU

¢n

=1¢

uJ3

' =- IDLE(5 MIN.)

I

1=, ii = i i i = i i m I i i i I ram=

TIME

Figure 90. MissioD Mix Cycle D.

, 164

®'i
I

ORIGIN,_.LPAGE [3;

OF POORQUALITY

Using these temperatures and corresponding stresses, the

stress-rupture damage at each cycle of the mission can be evaluated

for the critical elements of the blade and rim. The results of this

damage analysis are shown in Tabl= 65. This table shows the

stress-rupture damage at the critical life section of the airfoil,

(

approximately 30-percent span (Figure 91), for each mission cycle.

I

{" Damage analysis of the SC blade indicated that the maximum stressrupture damage point on the blade was at the "high c" point on the

i critical section (Figure 92). At this point, approximately

25 percent of the stress-rupture life of the blade was consumed

|

•;" during one 200-hour test cycle. Thus, the calculated blade life is

800 hours of mission life or 4 of the 200-_our test cycles at

w

. i096°C (2005°F) T 4. The calculated damage due to low-cycle •

_ _ fatigue was found to be negligible even when the fatigue life was

_•'_ based on DS Mar-M 247 material properties.

TABLE 65. STRESS-RUPTURE DAMAGE FRACTION.

: I _, rfoil , Blade Firtree D_sk Flrtree _:

I ._qh "c" I'Trail ' znglI '.

_ _ Cycle Point Edge IL Lobe 1 Lobe 2 Lobe 3 Lobe 1 Lobe 2 Lobe 3

=. A i 0.028 0.02_ <0.001 <0.001 <0.001 0.016 0.029 0.019 ,

0.Ii] 0.IIi _0.001 <0.001 <0.001 0.065 0.112 0.074 "_

i _

C I 0.048 0.044 <0.001 <0.001 <0.001 0.029 0.049 0.033 ; _ L" i

D I 0.060 0.056 <0.001 <0.001 <0.001 0.034 0.059 0.039 !

i . !

Total j 0.247 0.237 <0.001 <0.001 <0.001 0.144 0.249 0.165 k

: ..... I

/

, It

_' ¢ i

,_ Comparison of the predicted life of the SC turbzne blade to t

P

, the predicted life of oth cooled equiaxed and uncooled DS turbine ,_ '

blades shows that the SC blades offer a considerable advantage over '.

these more conventional materials for advanced high temperature "

:> operation. The normalized stress-rupture (S-R) life of the SC _,

" 165

i
. q)

ORIGINAL PAG'_ |5

OF POOR QUALITY

: Figure 92. Maximum Stress-Rupture Damage Point.

166

4'

,A
ORIGINAL FAC7 !;;

OF POOR QL',-"_::" '

NASAIR i00 turbine blade is compared to other turbine blades in

Table 66. As shown, the stress-rupture life of the SC blade was

more than twice that" of any other blade analyzed at the elevated

operating conditions. (Cooling air was held constant for the

cooled-blade analysis. )

[

!-

" TABLE 66. COMPARISON OF NORMALIZED STRESS-RUPTURE

_;. LIVES AT T 4 I096°C (2005°F1 AN{)

i N = NTO RPML

_.. Blade Blade Normal ized

" Material Configuration S-R Life

•_ ' SC NASAIR i00 Solid !.00 _

B .

_ DS Ma_-M 247 Solid 0.36 '_'

'i-.-::

'"" iN-t00 Solid 0 09

_: iN-100 Cooled 0.39

I

t .

,. 7.3.1.2 Blade Vibration Analysis

_ The vibratory response of the final blade design configuration

t, at engine operating conditions was calculated using the modified _/

" computer programs. Blade vibratory response is influenced by two :

factors: rotational speed and blade temperature. As rotational

speed increases, the blade is stiffened by the centrifugal foLce

:" field that tends to straighten the blade and increase the resonant

;"-_, frequencies. Conversely; as operating temperatures are increased, ' 1

_' the stiffness of the blade and the resonant frequencies are de- '

:-_' creased, due to the reduction in material modulas at higher temper- ' i

_; atures. '

i

,:

} <"

._ Figures 93 through 96 show the influence of rotational speed

._ on blade stiffness at room temperature. The magnitude of this

- effect varies with mode shape, with first flex (Figure 93) showing

,_ the greatest effect.

C

167

i
ORIGINAL PAGE Ig

OF POOR QUALITY

,, ,% ,

6

NORMAL MODE SHAPE

RPM = 0 RPM = 1.027 NTO ,'.r,

" FREQUENCY (HZ) = 4896 FREQUENCY (HZ) = 5157

"_ Figure 93. Single-Crystal Blade Vibration Analysis.

• First Flex at Ambient Temperature.

%

NORMAL MODE SHAP

%

, RPM = 0 RPM = 1.027 NTO

FREQUENCY (HZ) = 111478 FREQUENCY (HZ) = 10474

Figure 94. Single-Crystal Blade Vibration Analysis.

":- First Torsional at Ambient Temperature.

I

168 @ _
ORfG,IHAL P_ f.¢'.

OF POOR QUALITY

t

,--. NORMAL MODE SHAPE

i

RPM =0 RPM = 1.027 NTO

t ,",r;'

-'. FREQUENCY (HZ) = 12849 FREQUENCY (HZ) = 13024

¢,1-,.

:. Figure 95. Single-Crystal Blade Vibration Analysis.

•: Two Stripe at Ambient Temperature. '

_.

1 ) '

• i

• o NORMAL MODE SHAPE / o ,

,

RPM = 0 RPM = 1.027 NTO

, i FREQUENCY (HZ) = 20152 FREQUENCY (HZ) = 20249

Figure 96. Single-Crystal Blade Vibration Analysis.

%

First Complex at Ambient Temperature.

• f

169

! I
7.3.2 Aerodynamic Design

The aerodynamics design of both the SC turbine blade and companion vane are unchanged from the DS blade configuration designed

in MATE Project i. For aerodynamic details of these components

l

such as vector diagrams, aerodynamic loading, flow parameters,

etc., see NASA Report CR-159464 (Volume I).

}

! 7.3.3 Design Modification_

1

_|. A design modification of the blade tip was investigated to

...

improve the aerodynamic efficiency of the uncooled turbine blades.

This design incorporates a small tip winglet on both pressure and

: suction sides to reduce the aerodynamic losses due to flow across

_ the blade tip. To avoid aerodynamic interference, these tip winglets were added on top of the current DS blade design, thus making

,_ the SC blade slightly longer. This required a change to the ID of

_ the shrouds to maintain the current rotor-to-stator clearance and

i

thus avoid a blade tip rub. Stress analysis of the modified blade

indicates that the blade load is increased less than 2 percent with

the addition of these winglets, and the stress at the tip section

L

is only 27.6 MPa (4.0 ksi). These small changes in centrifugal

&

; stress will not appreciably change the mission life.

Design studies were also initiated to reduce the cooling air

w

flow in the rim area of the HP turbine. As a result of these

studies, an aft flow discourager was added to the blade with the

tooth extending radially inward. This design increases the load on

, the disk approximately 5 percent and introduces small additional

{

moments to the firtree. These additional forces have a negligible

impact on firtree life. A boundary layer dam (platform seal) was

)

also inserted between each blade to reduce platform leakage. This

sheet metal seal increases the blade load slightly more than 1

percent and has no significant effect on either blade or disk

: life.

170

®
ad I+J

ORIGINALPAGe._5

OF POORQUALITY

Figure 97 shows several views of the final design configuration of the SC blade that incorporates all the added performance

features. Two drawings of this blade were issued: Part Number

3076078 for SC Alloy NASAIR 100 and Part Number 3076086 for SC

Alloy 3. Only the material and heat treatments are different on

these drawings. The cast and machined configurations are unchanged.

E _

_ _ E

1 ""

|

s

!

I SHANKMODIFIEDTO

• \ ACCEPTPLATFORMSEALS

AFTFLOW

DISCOURAGER " i

>

Figure 97. MATE SC Blade - Final Machined Configuration.

• _ 171
SECTION VIII

8.0 COMPONENT MANUFACTURE

The objective of this phase of the program was to manufacture

and qualify at least two complete sets of SC turbine blades of the

new design for subsequent component and engine tests. This task

• included modification of the wax pattern tooling, casting and inspection of components, and machining of blade attachments and

tips. At the conclusion of the casting campaign, an economic

analysis for SC blades in production quantities was completed•

I

8 .i Tooling

r

,- Wax pattern tooling for the new MATE SC turbine blade was pro-

: duced by modifying the existing tooling from the Project 1 DS tur-

,. °

" bine blade. This approach saved both time and money, since the

.: basic airfoil shape was not changed. The modified wax pattern for _

the new MATE TFE731 HPT blade and the blade arrangement in the mold _

for the casting campaign are shown in Figure 98. ,_

o

8.2 _Casting Campaign _

Turbine blades required for component and engine tests were :_

%,

" produced using the technology described in Section 3.0. The '"

' NASAIR i00 material specification and acceptance standards for SC ; D

NASAIR 100 are provided in Appendixes B and C, respectively.

_ e

_' Jetshape's overall yields for the casting campaign are provid-

%,

¢ ed in Table 67. As shown, process yields were below expectations. :'

The results of these inspections reflect the problems that must be .:

_, overcome before the exothermic process is acceptable as a produc- ,,

tion process for single-crystal blade and vane castings Most of '_

.g

these problems can be corrected by automating the various facets of ,_

: the casting operation including mold handling, mold cleaning, exo- _i

173

r'._ pRECEDING PAGE BLANK NOT _ t
Figure 98. Wax PatteCn (Lower) and Assembly (Upper) for

2 Modified SC TFE731 HP Turbine _lade.

174

I I
cR_G_!_L P_G'z _

OF poOR QUAt'F[_(

TABLE 67. VENDOR INSPECTION SUMMARY FOR SC TFB731 BLADES.

u

Total Rough I Single Grain2/

Blade Visual Visual FPI X-Ray Laue Overall

: Material Cavities Accept Accept Accept Accept Accept Yield

Full Program Results

NASAIR I00 384 331 177 131 80 60 16%

ALLOY 3 288 263 176 134 99 74 26%

Final Six Molds Cast

• ASLOY 3 72 71 52 48 32 29 40%

;_ " 1 • Incomplete fill, hot tear, mold runout, cracking, gra|ns

2. Vendor evaluation of SC grain and overall visual acceptance

-: thermic packing, preheat and exothermic burn, transfer to the casting furnace, and final casting.

4 The processing at Jetshapes was done by production people, us-

; ing existing conventional casting equipment with minimal special-

,_ ized handling _quipment. Although they succeeded in surmounting

" some of the problems of the environment, the yields reflect several

ongoing problems in castin 9 these alloys. Improvements in processw

', ing are reflected in the improved yields shown in Table 67. These

were realized as production of the engine test parts progressed,

particularly in single-grain development and in overcoming initial

_ blade losses due to casting runouts.

The FPI and X-ray rejects were predominantly due to inclusions. FPI also revealed a problem with platform shrinkage on the

leading- and trailing-edge ends, indicating a need for feeders at

_ those locations. The greatest source of the inclusions came from a

- lack of mold cavity isolation during exothermic filling, handling

v

. 175

............. I _ _ _ .i !
in and out of mold firing, placement and movement in the vacuum

casting unit and, finally, crucible breakdown and refractory

spalling due to the high casting temperatures required for crystal

development.

l; Considerable progress was made in producing s&ngle-crystal

,.-

:. castings, but work must be continued on the exothermic process to

i reduce the pickup of ceramic particles and dross in the casting

process. Proper handling equipment, furnace shielding, and a

!: greater degree of mechanization for the production process will

,. eliminate many of these problems.

Comparison of the two alloys used in this task appears to inb dicate that SC Alloy 3 has a slight advantage in yield of SC blades

,v '_'

produced, while NASAIR 100 had somewhat less of a problem with excessive shrinkage. Since most of the NASAIR 100 parts were cast at

the beginning of the campaign, the processing experience gained

_ during this task was more beneficial to the Alloy 3 SC _iades.

The exothermic pack process is a viable production process for

SC castings and could be made better with more process/foundary

development work. Its major advantage is the capability to produce

: SC castings using only slightly modified existing equipment. Its

development into a competitive SC casting process for production is

an attainable goal with major additional equipment, such as a manipulator for mold handling, multiple chills and cooling sections,

and mechanical handling.

' 8.3 Casting Inspection and Testiqg

; Examination of Laue X-ray diffraction data indicated that me_,:t

of the blades were well within the orientation requirement ch£,t _._

[001] direction must be within 15 degrees of the sta3k_r_ a_is.

_ For the initial group of 102 blades examined, 95 percent _i[ bla,:::_

" were within i0 degrees of the [001] as shown in Figure 99.

176

,%

¢ ® p .._, ........... --_.,,..,nm,,,_aanlbw,_ w..,.- _ m • _l

in,
r>,_a I,,R

OF POOR QUALITy

0

1 2 3 4 5 6 7 8 9 10 I1 12 13 14 15

DEVIATIONOFSTACKINGAXiSFROM[001]INDEGREES

Figure 99. SC Blade Orientation (Laue) Results.

Mechanical property tests--760°C (1400°F) tensile and 982°C

(1800°F) stress rupture--from final design SC blade castings of

both NASAIR 100 and Alloy 3 exceeded specification minimums. These

data were included in Tables 40, 43, 47, and 50.

8.4 Blade Machining

After the blade castings were received at Garrett, they we:e

inspected using visual, radiographic, and FPI methods and shipped

to Walbar for final machining. The final machined blade configurar tion is shown in Figure 100. One hundred and thirty-four blades

177

.............. ,..: ..... m,.._- ,,m,,L _._ IF i
I

I

/

ORIGINALPAGe.IS

OF POrR QUALITY

Figure i00. Fully Machined MATE Single-Crystal Turbine Blade.

were machined: 74 NASAIR 100 blades and 60 Alloy 3 blades. None of

these blades were coated. Two blades were dimensionally discrepant

and subsequently scrapped. FPI of the machined blades again re-

_ %ealed the platform porosity previously noted on the blade casting, :

: but no other significant discrepancies were found.

'; 178
!

SECTION IX

9.0 COMPONENT TESTING t

if Due to the anisotropic material properties of single-crystals,

; there was concern as to the capabi._.ity of present analytical compui" ter programs to properly predict all vibratory modes. To avoid

problems during e_ ine testing, a two-phase component test program |_

_! was included _o correlate present analytical techniques with test

I:

! data and to confirm the SC blade design, i

t

' 9.1 Holographic Testing

b[ The first phase was holography testing ef eight SC NASATR i00

blades and four SC Alloy 3 blades. The blades were clamped at the

• z

" - root in a firtree type fixture that simulated turbine disk and seal

" plate loads. The clampe_ blades were excited to low levels of

:_ motion by a small crystal driver, while laser equipment produced

\ hoi_graphic interferograms to determine the natural frequer.cies and '

mode _hapes.

:m Resonant frequencies between 0 and 30 kHz were determined, and

• mode shapes were recorded both by still photographs and a real time .

' video recorder.

e

' Results of the holography testing indicated that the addition ,_

of the tip winglet lowered the natural frequencies of the blade

compared to the Project 1 DS blade, but did not significantly alter

the resonant mode shapes. A comparison of the first three natural

frequencies of the DS Project 1 bl_de and the SC Project 3 blade is

shown in F_.gure I01. A frequency diagram, shown in Figure 102, was

_ generated using the holography data corrected for material t_mper- '

ature and centrifugal stiffening.

179

@'

! 1 ' '
OF POORQUALITY

MODELTFE731 MATE

0S HPT61.ADE SC HPTB_DE

J . T "

k

• , _ . _41pt._:" " FIRST ._

FLEXURAL _ -

i .._,'_ _-,.

t_ 3400 _'_ 3198 HZNIO0

_. 3212 HZA3

, d

" ..... B _"-",,I ":

" FIRST _ ',

- , TORSIONAL :

f

2

% 8700HZ 7077 HZNIO0

7046HZA3

"" "_ t', /

_OR_IO_AL " "

"" .: , " COMPLEX .-

" i _

11,200H_ 9197HZNIO0

9407 HZA3

,.-°

t_!clure i01. Compariscn of Holography Results

between DS and SC Turbine Blades.

18f/

t •

'_' _
L

ORIGINALP;:,C;,}'t_

OF POOR QUALiT*'

Figure 102. M_TE Project 3 Single-C_'ystal HP Turbine Blade

hol_raphy Data and Operating Point Frequency

_'_" Prediction •

• 181

p

1

l

., , _ , ,t _'o

r
i

ORIGINALPAGE IS

OF POOR QUALITY

Stress-coat testing had been planned to determine the optimum

strain-gauge locations, but due to similarities between the resonant mode shapes of the Project 1 DS blade and the Project 3 SC

blade, stress-coat testing was not required. Strain-gauge locations were similar to those used iA, Project 1 componept testing.

£

,7

! 9.2 Firtree Testing

Jf

t The introduction of single-crystal castings with anisotropi _

properties makes life prediction of the firtree attachment area

'.;. very difficult. A comparative subcomponent firtree test of various

materials, therefore, appeared to be the fastest and most straight-

. forward method of evaluating the strength of single-crystal cast- _

5 ings in the f_rtree area. _,

A double-ended male firtree test specimen and corresponding

_ firtree specimen holder (shown in Figure 103) were designed with the

=!*

I

r _

•' , t '

•

• i

" 2, _ ?

m ¢

i

-" ) 1

P71_qNS-2

i

Figure 103. Firtree Specimen and Holder for

>

. Project 3 Turbine Blade.

182
same firtree configuration used for the Project 1 turbine

blade. Firtree specimens were machined from cast single-crystal

NASAIR i00 bars and from DS Mar-M 247, EQ MAR-M 247, and Waspaloy

bars for direct comparison of tensile, low-cycle-fatigue, and

( stress-rupture capabilities.

Testing was done in air at atmospheric pressure at a tempert ature of 760°C (1400°F). Results of the low-cycle-fatigue and ten- i

sile testing are shown in Figures 104 and 105.

Stress-rupture testing of the SC NASAIR i00 material was at-

% tempted at a stress level of 758 MPa (ii0 ksi) for an expected life _

of 500 hours. Unfortunately, a failure occurred in the Astroloy _

_- _ firtree specimen holder in less than i00 hours. Examination of the

failed holder and the SC specimen indicated that the test load at

_ 760°C (1400°F) was beyond the capabilities of the Astroloy firtree

holders. These results were revlewed with the NASA Program Mana- _,

: ger, and the remaining stress-rupture testing was c_nceled e _ t

An examination of the data collected during this testing indi-

; cated that the SC firtrees were at least as strong as firtrees made

, from the more conventional EQ and DS materials currently used in

Garrett production HP tu_In_ _iades. Therefore= it was concluded

that the Project ! firtree configuration was more than aeequate for w

: the MATE Project 3 blade design.

9.3 Rig Testin_

The second phase of the component tempting task determined the

I:

.._ vibratory response of the SC blades in an actual engine environment

using Garrett's high-pressure rotor rig (HRR). The HRR allows

strain-gauge testing of the high-pressure rotating group while

operating at actual engine conditions, i.e., 100-percent speed, .

- • maximum continuous turbine inlet temperature.

ct

183
ORIGINAl ;AGE _S

OF PO0:, QUALITY

1379(2001

WASPALOBY

1241llSOI

110311601

, 27614Ol 0 2601500l 538IlOOO) 816115001 1093120001

TEMPERATURE, °C (°FI

- Figure !04. Firtree Tensile Test Results.

" 184 J

i ®
ORIGINAL PAG_ tS

OFPOORQUALITY

172412501

137912001 % ' NOMINALWASPALOY9

E

t_ I,N1034(1501

t = -. r,_ C)

_" . _. .".r:

,.;. 689(ll_l

"" © SCNASAIR100 0

'" © 0S MAR.M247

:_ 0 EQMAR.M247

'_ 0 WASPALOYB

.:- 345150J.....

_02 103 to4 lob ,_

_" CYCLESTOFAIL

"_ Figure 105. Firtree LCF Test Results at 760°C

(1400°F) and R = 0.i.

• A

;' Twelve SC blades were instrumented with dynamic s_rain g_uges <

and placed in the rotcr assembly. Figures 106, 107, and 108 show, b_

respectively, the strain-gauge locations, the iqstrumented blades,

. and the completed rotor assembly containing the instrumented

blades.

"% HRR testing consisted of two slow accelerations and decelera

tions t_rom idle to maximum takeoff to idle, recording strain-gauge

output continually. During testing, a moderate resonance with a

z

peak-to-peak straln level of i000 /_inch/inch was noted at 21,300 ,:

_pm. No _ther significant :esponses were seen. The rig test was b

completed without incidence and the parts removed for routine post- _

" test inspection. '

, 185

i,
4'

!

ORIGINAL PAGE IS

OF POOR QUALITY

/t_CKING AXIS

• _ OFROTATION

f:

n. [] ,.o.

I -

• ._

36.1mm 0.25mm 40.6mm

_ 11.45IN] I0.01IN] TYP 11.60INI *

.,_-

,, I

• .].

: 23.6mm

•* 10.93IN) ": 16.1mm10.635IN) ,

i ! j,

,; -

-" _ 15.9mm_ 13.6

-" !- le._251N) _ w _ OFROTATION ..

" BLADENO. IOENTNO. DISC.IOC. POSITION,

A,

A59 $0 6402 12 LE

•_ A44 SO8402 44 LE

A26 $06402 26 LE

A20 SD6402 60 LE

N4 SO9872 4 TiP

A37 SO9e72 24 TiP

A57 SD7576 8 CENTER

; N24 SO757P 40 CENTER

A35 SO7508 46 TE

A41 SD7598 21 TE t

A12 SO2508 52 TE !

_,, N17 SD2598 32 TE

. Figure 106• Strain-Gauge Locatien, Single-Crystal

HP Turbine Blade.

: I_6 i
,,f ORIGINALPAGE i_

OF POORQUALITY

I:

!

!

.=-"
i<9 !

',, ORIGINAL PAGE Ig

OF POOR QUALITY

TRAILING

EDGE mb,.al,,._ _ a,,.. LEADING

f a,,. a,. EDGE TRAILING _ ,,_

EDGE . _, "" .

}" , .

. EDGE , ,6#" ' TIP

_ 4* • " % \ "

%

|, , ,_ • _ VJ

• _ _ _ CENTER

_: CENTERdi_F_ ' •_.' "_ '

"' LEADING "

-

[ _ . '_," EDGE _,

',_- _,

2_

Q

TRAIL:_G,

EDGE ,.. ,

, W P78530-5

.i LEADING TRAILING

;-' EDGE TiP EDGE "

:. Figure 108. MATE Project 3 Rotor Assembly Showing

• Locations of Instru_lented Blades.

t

9.4 Post-Test Ins pec toi.____nn i

!

I

J This inspection of the single-crystal blades unexpectedly i

revealed four SC Alloy 3 blades with leading-edge cracks at approx- !

_.mately the three-quarter span. No problems or abnormal conditions

were fo'md on any NASAIR i00 blades or any of the seals and shroud

segments. These four cracked blades were submitted for a complete

m_terial analysis.

. 188

®'
Results of the analysis indicated that the cracks all occurred

at the boundary between a recrystallized secondary grain and the

primary grain. These recrystallized secondary grains were all

approximately the same shape and size and were at the same location

on each of the cracked blades. Figure 109 shows a recrystallized

grain on one of the cracked blades (S/N A41) after macroetching.

_ Twinning was also found on the suction side surface. No signs of

fatigue were found on the fracture surfaces of any of the four SC

_ Alloy 3 blades.

i:

Further examination of _hese blades indicated that the den-

' dritic microsegregation pattern had been preserved during formation

of the recrystallized grain as shown in Figures ii0 and i±i. The _

recrystallized grain microstructure did differ from the rest of the J

_ blade in that no sub-boundary networks were present within the

recrystallized grain.

_i The fracture surfaces of the cracks in all four SC Alloy 3 ;' :

t ;

blades were quite smooth, curved, and lightly oxidized (as shown in ,:

r

Figure 112). The grain boundary cracking was attributed to a combination of centrifugal and bending stresses applied to the

extremely weak grain boundary that occurred during HRR testing. No

• indications were found on the fracture surface that were characteristic of a high-cycle-fatigue fracture.

" Recrysta!lized grains of this type are clearly unacceptable

. _nd attributable to an improper grain etch inspection after solution heat treatment. HRR test times and temperatures were not even I. p

close to the levels required to produce recryst_llization. I

Following analysis of the four cracked Alloy 3 blades and the "

_ discovery of extraneous grains at the crack site, all machined ,

L

blades and castings were thoroughly reinspected. This inspection in-

: cluded macroetching, a detailed visual review (40X magnification), ..

_'_ a sensitive FPI using Group 7 penetrant, and another radiographic

r _ 189
L
L
ORIGi_'IAL _ _*""" '_

OF POOR QUALIT"/

11 •

,; ,, f ,

z ' ,1_ ';.':

,_ ._- .ib _-

%•

i 1, •

•" 400_, . -"

'" I0'x0 , • /

E • . .,_00'_

,p' # i

; , _. :>-... .; --

2

• L

• ,,=1

L t[} X "-

]llU i , . . ,." , ; .

• I_111014

Figure iii. Microstructure of Recrystallized Grain from

Edge of SC Alloy 3 Blade S/N A25. Not_ Presence

of '_ins and Absence of Sub-Boundary Network in

, Recrystallized Grains. (Photographs are 83 p_rcent

of indicated magnifications.)

192
+,, !:i+

ORIGINALPAG_ i_

OFPOOR QUALITY
inspection. Extraneous grains and other problems were identified

in several additional machined blades and castings The results of

these inspections are shown below:

Rejected Due Rejected Due to Rejected Due

to Extraneous Freckling in to Major

_ Grains Critical Areas Porosity

;_ NASA_.r_.i00 16 12 I0

'. Alloy ._ 12" 2 1

*Includes the four Alloy 3 blades cracked during HRR testing

More than 20 percent of all blades inventoried were scrapped

b

" following re-inspection. An additional 16 percent were classified

as "marginal" quality. All remaining blades were approved for

• full-scale engine testing, which is described in Volume II of this

repot t.

194
.

_ SECTION X

_J

i0.0 COST ANALYSIS

" It is extremely difficult to project a fair and accurate comparison of turbine blade manufacturing costs for new materials

" and/or configurations. Differences in raw material costs, vendors,

2"

_ blade configurations, economic situations, lot sizes, and timing of

." the order all affect the cost of a particular blade. To make this

7 comparison more significant, Garrett historic manufacturing cost

' : data concerning a variety of turbine blades was examined and nor- }

malized. The relative costs shown in Table 68 represent an esti-

• mate of these normalized blade costs based on production size lots

-_ of small turbine blades.

Comparing the cost of SC castings produced by the exothermic

_:_ process and the withdrawal process is also difficult and can rea- _

_: listically only be done by a casting vendor qualified to produce

_ SC turbine blades. These suppliers have all made substantial _ ,_

_ investments in withdrawal equipment and have all gained consider- _

able experience in producing SC castings by the withdrawal process, i_

"_ Therefore, it is difficult for them to project the exothermic pro- _

cess to a comparable point on the learning curve and thus give an ,_

_ "apples-to-apples" comparison. Jetshapes has made this projection, ,_i

_ however, and feels that an SC blade could be produced for approxi-

"_ mately the same cost with either the exothermic or withdrawal pro- _

'_ cess (assuming the automated equipment were available and the same

• level of experience). ._

Table 68 shows that the cost of an SC casting is about twice _i

' that of a DS casting for a given configuration. This increased _i

cost is offset by the increased capabilities of the SC components.

An estimate of the benefits of the SC blades to the overall engine :_

can be calculated using a fully flexible engine cycle with variable

bypass and pressure ratio. Table 69 summarizes the results of a _

life-cycle-cost comparison of an uncooled SC blade to an uncooled

_._ 195
DS blade for a typical business jet application. (For more details ,

_, see Section 2.0 in Materials for Advanced Turbine Engines Pro£cam,

_ Volume i, Garrett Report 76-212445-1, January 10, 1977.) As shown

_ in this study, the advantages of SC blades are quite significant in

advanced engine designs that fully expoit their high temperature

capability.

c

,! TABLE 68. ESTIMATED RELATIVE COSTS OF SINGLE-CRYSTAL

: TURBINE BLADES.

!

" BLADE CASTING TOTAL* •

d

-- TYPE COST COST ,z

r EQ Cooled 1.0 1.0

;" DS Cooled 1.9 1.6

%.

_. SC Cooled 3.5 2.7

EQ Uncooled 0.3 0.4 _,

DS Uncooled 0.9 0.9

SC Uncooled 2.1 1.7

' *IncLudes machining, coating, handling, etc.

'4

196
?

_;_ TABLE 69. RESULTS OF THE LIFE-CYCLE-COST ANALYSIS OF THE

*'_ UNCOOLED SC HIGH-PRESSURE TURBINE BLADE AS

_/< COMPARED TO THE UNCOOLED DS TURBINE BLADE iN

'_ A FULLY oPTIMIZED ENGINE FOR A TYPICAL

:-%

_ COMMERCIAL BUSINESS JET MISSION.

_::" ATSFC = -1.53%

AW t = 0

,_ Scaling effect (BPR) -

%_W t = -6 67%

"._.

: Scaling effect (PR) -

%AW t = 0

Total AW t = -6.67% + 0 = -6.67%

ACost = +0.98%

; Scaling effect (BPR & PR) - same as W t

% Cost - -6.67%

_: Total ACost - +0.98 - 6.67 = -5.69%

ALife (TBO)* - 20%

* L

AReliability (MTBF)* = 30%

|

; ALCC l

TSFC Effect: -i. 571

Wt. Effect: -2.868

Cost Effect: -1.525

TBO Effect: -0.091

MTBF Effect: -0.019

Total _LCC% -6.074%

ALCC, MS 2095 !

*Increased llfe and reliability may be expected due to greater

integrity of a slngle-crystal blade (no grain boundaries).

197

®
SECTION XI

ii.0 CONCLUSIONS ,

5

i.

_. Casting Process

_" The low-capital-cost exothermic directional solidification

.... process was successfully modified to produce engine quality SC

?

_: turbine blades.

.:

:. Overall SC casting yields were not as high as expected (about

-o 20 percent). However, since the SC grain yield was reasonably high )

;

: (about 60 percent), a significant improvement in overall casting

[ yield would be anticipated with automation of the mold construction

_ and handling facets of the process.

_" Material Properti, es

SC Mar-M 247 stress-rupture capability was not significantly

. better than that of DS Mar-M 247.

t

SC derivative alloys, developed by removal of all (or most) of

• the grain boundary strengthening elements (Hf, B, Zr, C) from Mar-M _)

: :- 247, permit SC components to be optimally heat-treated and achieve ;:

: significantly improved stress-rupture capabilities.

/

Mechanical properties (creep-rupture, tensile, LCF, and HCF) .:

• of SC NASAIR i00 and SC Alloy 3 were significantly superior to .:

those of the DS Mar-M 247 parent alloy, with SC NASAIR i00 slightly :_

_ _:_ stronger than SC Alloy 3. ;_

SC NASAIR i00 exhibited evidence of mu phase and = tungsten in : _ the microstructure. However, this microstructural characteristic

_ did not produce any degradation in the alloy's mechanical proper6 P/_ECEDINGPAGE BLANK NOT FILMF/)
ties (stress rupture and HCF) even after long-term, high-temperature exposure.

Physical property data (density, thermal conductivity, and

i.

, thermal expansion) for SC NASAIR 100 were found to be equivalent to

; data previously obtained for DS Mar-M 247. Elastic modulus for SC

NASAIR 100 was slightly lower than published data for pure nickel

single crystals.

Retention of some Hf in these derivative SC alloys improves

fi. coated oxidation/hot corrosion resistance.

Blade Desi@n

i

:. The MATE Project 3 SC blade has significantly better fatigue

%.

and tensile capabilities and twice the stress-rupture l_fe of the

Project 1 DS blade.

Analytical techniques, modif led to accept anisotropic

. material properties, can correctly predict the lower mode vibrational response of the single-crystal blade.

Componen_t Testing

,4

No significant relationship between crystal orientation and

blade vibrational response could be determined for the range of

orientations examined.

Vibrational response (nodal shapes) of the Project 3 SC blade

was the same as that of the Project 1 DS blade. The addition of the

tip winglet to the SC blade, however, lowered the resonant frequencies.

Secondary grains with significant orientation mismatch cannot

" be tolerated in moderate-to-high stress areas.

200
l i%

[ •

i ,

_,' Ti_e Project 3 SC blade was 3udged acceptable for full-scale

, _L engine testing.

_ °i? Cost Analysis

i In production quantities, the cost of the Project 3 SC blade is

_, estimated to be about twice the cost of the Project i DS blade.

_ !_ Gener al

!. _

The increased temperature capability of the SC turbine blade :

_ will allow a significant increase in engine performance. This

._ capability can be used through increased TIT, decreased cooling air

<_- flow, and/or increased component life.

I 6

.ll

t

201
ORIGINALPAGE IS

OF POOR QUALITY

}

? APPENDIX A

:, MEASURED FREQUENCY RESPONSE OF PD SC TURBINE BLADES

¢ TABLE 70. MEASURED FREQUENCY RESPONSE OF PD SC TURBINE BLADES

,_" WITH VARIOUS CRYSTALLOGRAPHIC ORIENTATIONS.

|

" 8ecLal Crystal first i rLrst Torsional Second Two UnLdentLfied

NO. OEientation Flex Torsional Complex Torsional CoIp1ex St_ip_ NLsc. Nodes Nodes

-" 405-12 [240]-8" J 4085 8453 12540 24734 2650B-like 2T 19194

405-2 llOO]-I;'i 473i 7977 23.5 ii545 20,99-e 962o

°° r

408-12 [310]-8" 4097 8429 23363 26444 30365-3T 4065, 6136,

, i 10547, 16566,

i I 18721 20194

t ,, ,_

343-14 1100l-6" 1 4953 I 10131 23551 26233-(: 0570

l t I 31528-3T

". 397-8 [331]-6" i 4382 8519 _ 27105-C 26013, 30447

405-7 [100]-10" 4096 8471 24335 31424

. J .....

407-1 1113l-9* 4192 8513 17544 24973 17994-C 17191, 17357

_" r 2till-like 2T 17517, 18279

26189-,C 25211, 27597 " ':

28180

r 343-10 [210]-9" 4776 8273 27124-¢ 18800, 19126

- 343-_" [I001-23" 3717 8085 ; 13090 17577 -- 31720 4475*Tip Flex 6847, 28552, 1124J-Tip flip 11100

23663

I/

394-3 1100]-5" 4806 9441 lilit 22692 29000 t.2l, 6.21,

i ll.2K, _).lX,

I 12.IX, 26.3K, F

23.9K, 28.9K

..... I

: 3iS-i [210l-i" 4502 977I i 12671 19723" 23062 203il 17.21, 29.21-° l.il, 11.21,

; _ 25.41

." 307-5 |100]-5" 4375 9593 i 13092 19685 24094 17.5K, 28o4_ 12olK, 13.5K,

I 29.2K

387-3 i100]-5 e 4514 9446 13026 19221 12.6K 12.1K, 19.6K

l II 21.3K. i3.1X,

24.21, 26.11,

i 29.91

: 387-1 [100]-5' 4537 9969 12372 I 11713 22964 24290 /ll 2756,17K, 25.7K5104'

.... -- # " ' I

1344-7& [10_l-2.2" -- 9761 12427 1 20114 23027 "" I 31304

i 343-51 [1001-23" I 4370 9096 13101 17117 .... * 19453, 24454 i

• t lii71 I

f 343-15C 11101-11" 4189 7920 12512 1 14412 21511 22707 , 20004, 34066

343-13D ll001-I" 4105 8515 11659 ! 19230 21115 i3923 * 33723

..... l ..

323-51 11001_4" 3841 8826 .... 21322 25110 27911 l

,- 30514

, ,] 343-14r 11101-16" 5633 9701 15583 19481 -- 22321 216t[, 24411,

32721, 28601o

33635

• i

L

L

202
b

ORIGINALPAGE15

OF POOR QUALrrY

APPENDIX B

3 NASAIR I00 MATERIAL SPECIFICATION

! ,NT I '

'MATERIAL Jl SPECIFICATIO i N 199!93 ] EM.,.1 ! .c

1. APPLICATION 3.3 Castxngs shall be poured only from

remelted maste, heat metal.

I.i This speclflcat_on establlshes the

•,_ _ requirements for NASAIR-100. 3.4 Blades shall be cast from every mastel

heat and tested per EMS52501,

' i.i.i NASAIR-100 xs a cast nlckel-base

• _ superalloy used for single-crystal (SC)

turblne vanes and blaues at temperatures 3,4.1 If the Configuration permlts, tes_

'i " up to 2000"}. specimens shal also be mach!ned from cast

parts, f

•' 2. APPLICABLE DOCUMENTS %

3.4.1.1 Speclmens may be machlned f_om any

2.1 The follow_ng documents form a part area of the casting, unless otherwlse

I. of this speclflcation to the extent refer- specified.

enced herein.

3.4.2 Blades shall be cast into the same

2.1.1 Gdrrett Speclflcat_ons type of refractory mold as the castlngs

for which the master heat is to be used.

" EM552300 Classlflcat_on and Inspection of

Castlnas ._ , ?

_. / 3.5 Castings shall be supplied in the

E_:552330 Master Heat Preparation of solution heat treated condltlon per EMS5250:

_" _ Nickel-Base Alloys

" _. 4. PROCESS CONTROL

EM_52501 Metallurglcal Control of

_ Directlonally Solidlf_ed Single- Not applicable.

Crystal NASAZR-100 Turbine

Vanes and Blades i S. 5. ZNSPECTION

t: ;" | "

2.1.2 Aerospace Materlals Spec_f:catlon 5.1 All castings shall be vlsually, pene- ! -"

trant , and X-ray Inspected i_ accorddnce

• AM_ 2280 Trace Element Control, Nickel with _52300.

-- { A_loy Castings

, 5.2 The suppller shall perferm all testln_i !

3. TECHNICAL R_OUIREMENTS for conformance to chemical l_mlt_.

3.1 Composition Suggested A_m Ran@e 5.3 The supplier _hal] perform all

: mechanlcal-property testing.

• _ Carbon .0.006 0.006 max.

Chromium 9.00 8.50-9.50 6. ZDENTIFICATZON AND PACKIN_

_ Molybdenum 1.00 0.80-],20

e

Tant_]um ].]0 ].10-_.50 _:gt applicable.

Aluminum 5.8 5.60-5.90

_,, Tltan_um 1.2 1.00-1.40

{ Boron <0.002 0.002 max,

Zirconium <0.03 0.03 max. 7. APPROVAL OR PROCU_;_NT

" Cobalt <0.50 0.50 max.

Tungsten 10.5 10.00-11.00 7.1 To assure un_formlty of guallty, sample

Manganese 0.20 max. castln_s from new or reworked toollnq %

Sulfur 0.015 max. shall be approved by the purchaser. S_Izcon 0.12 max.

Iron 0.I0 max. 7.2 Supplier shall use the same castsnq q

Nzckel Remainder Remainder technique, *nclud_nq rate of cool_nq after

'_ (O) _I0 ppm I0 ppm max. casting, and If heat treatment Is s_,eclfled,

(N) "12 ppm 12 ppm max. the same heat-trentln_ procedure for production castings as for approved salt 1(, i

3.]I Trace elements shall be controlled In castlngs.

-_ ac, rd_nce with AMS 2280, Class 2.

3.2 Production of master heats, remeltin_ _

of masttr heats, and pouring of castlngs

shall be accompl_shed under vacuum

3.2.1 Master heats shall be made tram 100%

vlrg_n mater_al in accordance with EMS52330.

3.2.2 A master heat xs defined aa previously

z efxned metal from a single furnace charge. 1

;

203

......
I

,b

oRIGINAL PAGES

OF POORQUALITY

|COO( IO(NT NO.|_."_..CIFICATiO_ NO. ! R(V LTR MATERIALI SPECIFICATION199193 ! II I .c

J

J' 8. l_pOR?S 9. OUJ_.,I'rY CON'I'POL

8.1 The supplier of c•stings sh•ll furnish 9.1 Castings shell be uniform in quality

[_ with each shtp_ent a report listing the and condition, lound, and fr•• from foreign <

results of the mech•nic•l-property tests, materiels end from i;,::rnel end extern•l

results of the chemical •n•lysis, end • imperfections in excess of those allowed in

statement that the castings conform to the EHS52300 for the specific class and grade.

requirements of this specification°

9.2 At the option of Garrett, • casting

• 8.1.1 This report shall include the put- shell be selected from any castings received

chase order number, master heat number and and shall be inspected in accordance with

_ code symbol, If used, material specification the •pplicable requirements for that part.

number and its revision latter, part number,

; and quantity from each heat. 9.3 Parts and material not conforming to

the requirements of this specification shall

=. 8.2 The supplier of finished or eemiftn/shed be rejected.

p•rts shall furnish wlth each shipment •

report showing the purchase order number,

materials specxflcation number, contractor

Or other direct supplier of castings, part

, _umber, and quantity.

, 8.2.1 When c•stings for making finished or ._

"- semifinlshed parts are produced or purchased

* by the p•rts supplier, the parts suppl_er

-: shall inspect castings from each master heat

or master heat lot represented and shall

include In the report a statement that the

castings conform, or shall include copies

of laboratory reports showing the results

of tests to determine conformance.

°

I III

204

..........®
' ORIG',HAL PAQE 15

1 OF POOR QUALII"Y

i APPENDIX C

:. ACCEPTANCE STANDARDS FOR SINGLE-CRYSTAL TURBINE BLADES

ii i i i ii

' MATERIALSPECIFICATION .o. :.* I 99193 i ;....1 1,,VNc,..,

L ,

L. APPLICATION 3.2 Master Heat Qualification

I.l This specification altablilhei the 3.2.1 Mamelting of master heats end pouring

-'. ' requirements for metallurgical control Df Of cestlnqs shall be accomplished under

turbine vanes and blades cast from vacuum.

NASAIR-100.

, 3.2.2 SC turbine blades shall be cast from

1.l.l NASAZR-100 IS I directionllly each master heat and tested. Teatlng shall

solid, fled s,nqle-crystal (SO) nickel-base consist of three 1400"r tenlilea and three

superalloy used for said turbine components 1800"r streii-rupture tests.

",' at temperatures up to _00"F.

3.2.3 SC turbine blades shall be cast into

1.1.2 When cast as a SC there is 8 signif- the swBe type of refractory mold as the cast-

:" it&me improvement in creep-rupture proper- ings for vhlch the master heat is to be used.

ties compared to multicolmmar-qrain

directiona)ly solidified (Da] and convert- 3.2.4 If the c8oting confiquration does

tiohally cast Jtickol-baie alloys, not alloy teat specimens to be machined from

" It, then Iample blade castings shall be cast

2. APFLZCAB wr DOCUN_TS e|onq trith the other cljtlnqa and test

," specimens shall be machined from these se_le

". .- 2.1 The following docuamnts foru a part blades.

• of thai specificatiol,, tO the extent refer- _f.

- , & Shoed herein. 3.3 Machenical-Property Teatinq

i

• _" 2.1.1 Garrett Specifications end FO_ 3.3.1 SC castings used for testing to

• qualify the salter heat may be qrit blasted

" I_52300 Classification and Inspection of prior to heat treatment provided the machln-

-' Castings ing of test specimens will remove sufficient

_'_' - material so ss to clean-up all cast surfaces

EM.S52309 Fluorescent Penetrant InIpeCtion

3.3.2 SC turbine blade cast,rigs selected .

£_t_52330 Master Heat Preparatxon of for mechanical-property testing ihall meet

Nickel-Base Alloys the followinq crystal orientation require- _ ""

smntl

; EMS52332 Metallurgical Fixed Process

Approval for Cast Turbxne Blade, |001] direction w_thxn 15" of the

Blades (Nochanical PropertieI) blade stacklna axis.

£M55234| _dioqraphic InSpection V_-rey Laue method is recommended for verlf*-

cat,on of grain orientation. Use of alter-

_ E15S$5461 Nickel Alloy Castings, Inveit- sate methods for 9rain orientation measurement, Corrosion- end |,'eat- gent shall require written approval of

J_llstent, NASAIB-100 Materials EnqineeTlnq. t

t

• i

Form P-53i9 Oual_ty ¢ontrol-Maquest for 3,3.3 Prior to test specimen machining, SC

Itateri81 Review Action (P)UtA) caltinq8 shall be heat treated as follows:

-- , Form P-0041 Fixed Frocesi Form 1} Solution heat treatment - 2400 _-I0"F for

two hoers, re*as temperature to 2415 +0

3. TECHNICAL REOUIKKIaENTS -10"F, and hold for two hours follo_d

; by rapid argon-fan quench to balow II00"F

3.1 Com_eitio_ The co-ling rate from the IolutIon heat

treat te_rsture must be sufficiently

].1.1 The chemcal colNposition shill be in flit tO meet mechanical propert,, requital _ accordance trltb B_J554SI. maStS. '

).1._ A master heat shall be made from 100t 2) Pseudo coating cycle he_t _reatmen_ -

_ - vlrgxn material _n sccord..,_-e with EMS52330. 1100 t_5"F *or five ho_s and still air

' _ 3.1.3 A maIter heat is defined es pre_ious-

_ ly refined metal from a sinqle furnace 3) Age heat treatment - 1600 _-_5°F for 20

t chsrqe, hours end alr cool,

l " '2O5
ORIGINAL PAGE li

OF POOR QUALITY

MATERIAL .... SPECIFICA _. T_ 'i_00(I_(NT_0.i_',l£,'T_ .0_ - t_v LTnI 99193 I ":

. 3.3.4 For mechanical-property testing, 3.5.4 l_talloqraphlc Inspection

tea_ specimen conflgurat£on shall bQ sa

follows: 3.5.4.1 A blade from each solution heat

croat lot shall be IRetallograpflica_ly exu-

,, Blade - Teat specimens sachined froR blades tned for Selma prime aolutionin9 end

shall have a O.070-1_ch diumter incipient silting. Acceptable and rejectable

qauge section 0,375 inch long be- _Lcrostructures (100X meqnification) are

trees radii, provided Ln Fiqursa 2 and 3 respectively,

A 100X pho*.o_Lcroqraph of I representative

Test specimen orisntation shall be area taken at the blede -_oot shell be subparallel to the stacking axis of l_Ltted to Garrote Roceivinq Inspection for

'.. the blade. Fig_tre I provides • transmittal to Materiels Enqineerinq.

•; Sketch Shovlnq the location of the

test spectJn on a blade casting 3.5.5 Surface Condition

" and a drawlnq of the teat speci-

:. man. 3.5.5.1 Production IC blade castings shall

be solution beat-treated prior to any

3.3.5 Tensile Testing - Tensile test spool- abrasive blaatlng siltation after tsamval of

.. ImnS shill be tested at 1400:10"F. Three castings from the wld. "

, spotless/ shall be tested with each storing

the following maises: 3.5.5.2 |lade surfaces shill show no .Jvidehce of recrystallisation, alloy depletion

L Ultl_te tensile strength (ksi) 173 or oxidation.

0.2 percent yield strength (_si) 146

- |longatxnn (percent in 4D) 3.0 3.5.1 Grain Orientation

-_. itQdl_cticm in ares (percent in 4D) 3.0

3.$,t.1 Each blade shall be chemically

; 3.3.6 Streas-_pture Testing - _trels etched for i time sufficient to lightly

t rupture teat specimens shall be teated under reveal the grain orientation,

I constant Itrlie of ]5,000 psi It llOO _5**lr. ,

Three apeci_na shall be teated and each 3.5.6.2 The |001| crystallographic qrowth

specimen shall have a mnimum li_e of 47 direction shall be within 15* of the |tackin,

houri, axis of the airfoil. Crystallographic orientation of all blue cast_ngs shall be deter3.4 Oualifled Master Heat mined. The orientation for each blade casting shall be docuJented and the results sub3.4.1 JQstsr beat _ter_al meetlno mechan- elated to Garrote Receiving Inspection for

• anal-property _Lnlmuaul specified herein is transmittal to Materials Engineering for

defined as a qualified m_ster heat. information purposes only. X-ray _lue ithod

is recmmended for qra_n Orientation.

: 3.S Ma:ellurgLcal Control of SC Turbine

V_nes and Blades 3.$.6.3 )io eq_i_ed _rains, col_er qre_ns,

or freckles are permitted on the blade east3.5.1 Caetinqe shall be po,_red only from ingi.

remeltsd qualified master heat _tel. The

use of gates, aprueI, r;sers, or rejected

' castings is not pe_ttl4.

3.5.2 Each solution heat treat lo_ met be

- - qualified.

3.5.2.; A solution heat treat lot Is da_ined

IS a Niter heat/pert _tuJl_Der/lOlUtion heat

treat _ycle combination.

, 3.$.] The solution heat :rest lot qualifit cation _y be performed I,_ conjunction with

{ the :Miter heat _uallficatLon.

, _ 3.5.3.1 All requirements for master heat

qual_flcat-on s_d _chanscai-Froperty teat-

' lag apply to sOlUtiOn heat treat lot qualification.

r

• , ,, , _

i

206

¢i ®'!
|

ORIGINAL i_AG£IS

OF POORQUALITY

]MATERIALSPECIFICATIONJ=91934! ' J

10(NT NO. SRICIFICATION NO. REV LTR

52501 N:

i

• I 4. PROCESS cONTROL . . I the supplier has made an inadvertent

| chanqe in the flxe_ process during the pro-

"4.1 Cooling rate fr_t solution-heat-treat duction of cas.,ngs, Garrett shall be imutditemperature oral1 be sufficiently feat to ately notified of the nature and extent of

: met nochancial properties, the change and quantity of parts affected,

and the supplier shall obtain vritten per-

_ 4.2 Solution heat*treat furnaces J_._ll be sissies via the serum1 _ procedure from

qualified by the cesti_ supplier. Gatratt prior to subu,_tting the parts re,

_. :eptance.

• 4.2.1 To qvality a furnace, the castit,g

supplier must heat treat a minimum of 5 4.4.3 llach sur, plier shall keep daily records,

blades per ear h test condition of tensile in accordance vith established procedures, of

testing end stJrese rupture test.tag epeci- 811 con,.re1 Atum An the fixed process that

fAod herein An a f_rnace loaded to a ere recorded continuously on strap charts or

"-" maximum production heat treat capacity, circular cherts, or checked and recorded by

and test to the scenarises1 property re- hand. A11 process controls shall be pers_nalquAremmnts. A sLmalated load by yoight ly audited by the supplier's q_ality -.ontrol

may he _.ted. aectAoc oJ_e_th, and a report filed.

4.3 All casting processes, including 5. XNaPECTXOII

natal amd mold prell_ratAon amthods, shell

'_ • be established and do_mm,,_ed, includin, j S.l |ub:)ect parts shall be inspected by the

controls for critical variables. YaAlure casting supplier vLth the relieving _'

:'. to meet the specified requirements shall mathodsl vLsual° radiographic, macro acid

= be cs_e for vit_holdinq qualification of etch, flucrescent-penetrant, sad dimensional,

, ' the castinq process. Gazrott revAev of as appl/cabAe.

;_ the process documentation and casting

evaluation shall ¢t_stitute q_alifAcatAon 5.2 Visual Xns_sctAon - kll bludel shall be

of the casting process, inspected in accordance vAth EMS52300.

: 4.4 &ft-r the casting process for each 5.3 Fluorescent Penetrant Xnspection

part number has been established, the

supplier snail prepare documentation, on 5.3.1 &11 blades shall be processed per

a Garrett supplied rased process Pore, of rJIS230J.

the final detailed feundary operations end

processing sequences. ThAn documentation, 5.$.2 Fluorescent penetrant indications

_. comprising the "fixed process," shall he shall he :_rrelatod v_th the allovable

s_JDl_ttted to Garrote for revte_, pr_or to vis_al imperfections and the eccept/relo¢_

the fLrs_ order of pruductzon parts, criteria of FJSJS;t300 as applicable.

4.4._ Xf At becol_s necessary for the

suppller to _e any cheese An this fixed

process, vrAttea peZlASSioa shall be obLaLned fro_ Garrett pr_or to incorporating 8vch

• chdmge.

:, _

! [

1

! :

ii

r t

f ;

207
!

;' ORIGINALPAGEi9

OF pOOR QUALITY

I I Ill I II I

ICOOE IOENT NO.ISPECIRCATIOfl NO. I REV £TR I

IMATERIALSPECIFICATION99193 E. s.O1 .c

t

5.3.3 Evaluation of smeared or uneharp 6. IDENTIFICATION AND PACKING

indications may be performed by wiping tn_

indication one time only with a swab or Not applicable.

! brush dipped in solvent.

5.4 Radiographic Inspection - All blades

shall be radiograFhically inspected per 7. APPROVAL OR PROCUREMENT

'_ " EHS52348 and the acceptance standards

deflned in EMS52300 as applicable. 7.1 Approval of the supplier's fixed process

_" and process changes shall be in accordance

5.5 Master heats shall be tested by the with EMS52332. All subsequen: fixed process

casting supplier for conformance to changes shall be approved by Garrett.

." chemical limits. Chemical analysis shall

be performed on a blade cast from the ft. REPORTS

|! master heat and meet the requirements of

EMS55461 B.1 The supplier of the castings shall furnish to Garrett Receiving Inspection with

5.5.1 Overall chemistry _-y be deter- each shipment a report listing the results I

i. mined at any location within the blade, of the mechanical property tests for each

part n_u_ber/master heatsolution heat treat

5.6 The casting supplier shall test mix lot comblnation, the results of t_.e chemical

blades from each part number/_a|ter-heat/ analysis from one casting per master heat

solution-heat-treat lot combination, three in representing the part number shipped, number "_

elevated temperature tensile and three in of pieces per master heat/solution heat ._'

creep-rupture per the conditions of treat lot combination and statement that

tensile testing and s£ress-ru_tuze testing the castings conform to the requirements of ;

c Specified herein, tc verify conformance tO this Specificazton.

_, the mechsalcal-property requirements.

" 8.1.1 This report shall include the put- !

_- 5.6.1 The specimens shall meet the mechani- chase order number, master heat number and

I, cal property reauirmnents of tensile testing code sTnbol, if ua£d, solution-heat-treat

and stress-rupture testing specified number, material specif_catlon number and It_

_ herein, reviszon letter, part number, class and

_" quantity from each heat. "_'"

5.6.1.1 If the mechanical properties do

not meet the minimum requirements, the heat 8.2 The supplier of finished or semifinishe_

treat lot shall be retested using two more parts shall furnish wzth each shipment a

blades for each failure from the same master report showing the purchase order number,

._ heat-heat treat lot combination. Any materials specification number, contractor

failure of a retest rejects the heat treat or other direct supplier of castings, part

lot. number, class and quantity. _,.

5.7 A sample from each heat-treat lot 8.2.1 When castlnas for making flnished or

received shall be inspected by Garrett semifinished parts are produced or purchased

for recrystalllzation, alloy depletion, by the parts supplier, the p_rts supplier

and oxidation, shall inspect castings from each master heat

or master heat lot r_pres_nted and shall

include in the report a statement that t_,e

_. castings confocm, or shall include copies of

'laboratory reports showing the results of

tests to determine conformance.

D

!

! I e IS i

" I

r

• , 208
®

r-

' ORIGINALPAGE IS

i OF POORQUALITY

I i

:. . ... 19 9193 I _.5o_ ,_

!_ 9. 0UALITY CONTROL

9.1 Castlngs with indications of shrinkage zn excess of Garrett Specification

EMS52300 requlrements shall be re]ected.

" 9.1.1 Castings with linear indxcations

• shall be rejected.

9.2 Garrett shall accumulate production

•. " audit data for Informatlon only, Five

blades per 5000 blades received shall

i'i be Jubmitted tO the Materiale Laboratory

l, on a CNR to study the effect of fin_shing on the cast blades. The blades shall

be traceable by master heat number/

solution heat treat _ot combination.

9.3 Parts and material not conforming to

the requlrements of this specification

shall be rejected. [

P

| !

I(

i :

E ,

' 209
!

ORIGINAL PAGE II

OF POORQUALIW

-- ,,i I I I I --

tMATERIAL ...... SPECIFICATION , , ,, ' i, 9=93, IEMSszso ,., z, . -- "¢ - .

±.

U-- ;_stackin_ _is

_ ) ',i,,_/_adin'ed''

_ I I_ /LongitudinaZZy

n; / j _/ Oriented 70-Mil Bar

! _ _ (Parallel to Stacking

:- _ Sis) t "

Sketch showing location of test specimen on

;" blade _st_g.

__0.071 5 - 40 Threads

• Typ0"75 R7 [-- 0.069 /

"

,_ -q:L o.,,, Jq I°''''

I 1.00 _l

; 70-Mil Test Specimen Drawing.

"L

II I I

Figure 118. Location of Test Specimen and

Test Specimen Drawing.

!

! 210
ORIGINAL PAGE ig

z OF POOR QUALITY

EMS52501

, Rev. NC ,'. KI

I_1 U r"

m 0 _ o

U •

0 0 ,-4 .I..I

'< 0 '

' i •;" 0 0 0 r_

.,-I

" _ _ '13 ,<.

- - ! _ c

, O_ _1

.. • - ,,,4 I "

, u _ _ o _ ._ _i,_ ,, [ r.J 4a

H _

• • .,'_E • _-,-I

4

• .,-_ I:: (J _0 -

_; ,

•,_ • ,-4

_ _ ._ ¢) •

r _ _ " 0 0 O _ 2_

, _ _ ._.I [

oo o _ !

, 211
!

l

_ ORIGINAl- pAGE ig EMS52501

,' OF PoOR QUALi'I'_ Rev. NC

/

• _ -,4

7" o_

" S -,- I_ _ rO

-_. 1 ,_ _ • _ .pl

.. rO .,4 O

0

;- ,. "_ ,,_

,_..._'_' ..... . ' .'_ 4J - " __;.;_,. _;_:.:..,:,_.':. " _, ,'._, . _11_ _b_lr' w " __ .

C_

"_ OO

- .,-4.,-I

: :3 _1

O.N

!

," ,e"'_' _,_ :. I_ _ ,-4 _ _E

.,-I

U_

_'_ .

_ 0 0

._ "% _ 0 _1 0

-, O O O

212
ORIGINALPAGE18 "

OF POOR QUALITY

i ii i • i i i

_ |CODE IL)LNTNO.! SI'_.CIFICATIONPK). I REV LTR

MATERIAL SPECIFICATION J 99193 I ,_52501(Appendixz, I NC

_._ ACID ETCHING ME_IODS

L

_ This appendix offers alternate methods for etching cast parts prior to inspection, These

_ ettochiexpose ng methgraoidns •bound re uari tileizes d or todendr accoitesmplishpriortwotopurmpacrogra oses, in (1) inspec to tioonbta, in anandetch(2) autofficobiintat in

:; • cleaning etch. When specified, the cleaning etch shall be used crier to fluorescpnt-

: , penetrant inspection.

CAUTION, Nixing of solution• end etching of parts must be accomplished in an area with

;' adequate exhaust ventilation, as toxic fumes ere liberated from the etchanta.

:, ;4e thod I /:

etching Solution:

100 gel Approx i liter 4

_'. Huriatic Acid (20" Be} 80 gel 757 ml ;

Anhydrous Ferric chloride, FeCl_ 135 lbe 154 g

Nitric Acid (42 e Be) 2 gel lg ml

"",_ Water 11 gel 106 ml

b 1 Add ferric chloride to murlatic acid. Allow to dissolve.

• : 2 Add nitric acid.

_._, 3 Add water. _ ._

_" a) A new solution shall be prepared when • suitable etch za not obtained within :

-,_ 12 minutes.

-_-_+" ' b) Do not raplenilh to maintain volume.

; Procedure:

": i. Load parts in etching basket, keeping level below basket rim.

2. Immerse parts basket in etching solution maintalned at room temperature (75-I00'F).

: 3. Check progress of etch after 6 minutes and every 2 minutes thereafter by removing

one casting, rinsing, and visually inspecting progress of etch. Once the etch

. " time required i8 established _or that particular run of castings, the following

loads can be run without check/hg. Typical etching tlme iS 6-10 minutes. @,

" a) Immersxon time for clianin_ etch shall be 20-30 seconds. _)

,- ? 4. Remove from etching solution and rinse in clean, cold water. '_

5. Immerse in alkaline cleaner solutlon for 3 minutes.

6. Remove from cleaner end rinse in clean, cold water. _.

. 7 Air-water nozzle scrub each individual casting clean.

e Blow loaded basket free of excess water with air only.

7

.;. Nethod_2

" __. Etching Solution : -._

Approx 2 liters

Murzatic acid (20' Be) 90% by vol (1615 ml)

'_, Glacial acetic acid 5% by vol. (B5 ml)

Nitric acid (42" Be) 5% by vol. t85 ml)

Ferric chloride to saturation (12.5 Ibs) q

I. Add acetic acid to muriatic acid while cautiously agitating the mixture, t_

2. Gently heat the mixture and add auf._icient ferric chloride to raise the boiling t

point to 150°160"F.

3. Cool saturated solution to_lOO*F, then cautiously add nitric acid while agitating

the etchant. C_UTIONt Never add nitric acid to the etchant when temperature is

above 100'r. ;_

a) The etchant shall be discarded when the etching time requires more than two

minutes to delineate the macro_rain structure. _

Procedure : 1! 1. Load parts in lUltebll tray or b_akat lo that elrfoill do not come in contact with

each other. _)_

213
®

!

ORIGINALPAGEIS

OF pOORQUALITY

iCOD_ IO_NT NO.iSPf.CaF_T_ NO. i REY LTR

MAI I TERIALSPECIFICAT | I I I IONJ99193 IJ .,5ol ..... I .c

}-. 2. Tnb_erle xn •csd etchant (150 _+10"f) for a minimum length of time to bring macrogreen structure viixble to unaided eye. N•x£mum exposure tame ;n the etchsnt

:' shall be limited to two minutes. The etchant or parts •hall be •g;tated to •ld

An obtainxr_ uniforD etching and to minimize the eXpOlUre tiM.

i) Immersion tJJue for cleaning etch •hell l-e 10-20 second•.

3, Rinse thoroughly In running t•p water.

C 4. Dessut by JJUmersing in concentr•ted hydrogen peroxide (HgO_, 35 percent). Hand brush or ' - air water power flush surf•tea of the etched i_tJ to remove residual

smut.

S. Rinse In running tap w•ter.

• 6. Rinse in hot llp water and dry.

_'• _'_._t

Etching Solution:

_o

Nur£etIc acid (20" Be) 90'& by vol.

Hydrogen Peroxide (30-3M_) 10% by vol. (or euffLcient qulntIty

to obtain 8 _tbtactory

*" etch)

b

L l. Add hydrogen peroxide to autistic acid while c•ut£ou•ly agitating the mixture. ,_h

a) Hake up solution _ust prior to usage.

._.. b) Whenever possible, the etching solution cont•iner should be immersed In a rap

water rinse tank for the purpOse of d;ssIpit£r_ the he•t liberated during the

'" etching process, 8o thet an etching tiue cycle can be established.

:; c) The etch•st shell be disc•tried when the etchang tL_e requires more thin five

; minutes to delLne•ta the macrogrein structure.

_: Procedure

1. Lo•d ixtrts In sult•ble trey or _asket •o thee airfoils do not oosm In contact vith

each ot_,er.

2. lnmersu in •cid etchent mk£ht_kned at room temperature (?S-IO0"F) for • minimum

lengr.n of time (5 sin. max.} to bring macrogrein structure visible to unaided eye

"- vhe:, inspecting for grain size and ciStiflg irregular/ties.

e) ImmerJion tJJmt for cleaning etch •hall be 10-25 8econde.

; 3. _/nse In running tap water. Hand brushing or air-water power flushing may be

: required if residual enut te not removed during the rinse cycle.

4. Rinse In hot tap water end dry.

t

I I II II I I

•

i

214
! REFERENCES

I. Sink, L. W., G. S. Hop_in, III, M. Fujii: Low-Cost

+_ Directionally-Solidified Turbine Blades - Volume I, NASA

it' CR-159464, January 1979.

-_. 2. Gell, M • , D • N. Duhl, and A. F. Giamei- • "The Development of

J_

' Single Crystal Superalloy Turbine Blades," in Superalloys

1980 - Proceedings of the Fourth International Symposium on

ii.:

_:" Superalloys, Champion, Pennsylvania, September 1980, pp• 205-

• _ 214.

a .

i

1 3. Barrett, C. A. and C E. Lowell: "Resistance of NiCrAI Alloys

:_+ to Cyclic Oxidation at Ii00 ° and 1200°C, ''Oxidation of Metals,

ii, 1977, pp. 199-223.

4. MacKay, R. A., R. L. Dreshfield, and R. D Maier: "Anisotropy

_ " in Superalloys 3

_i_: of Nickel-Base Superalloy Single Crystals, _ +_

+'_ 1980 - Proceedings of the Fourth International Symposium on _ -

i+ Superalloys, Champion, Pennsylvania, September 1980, pp. 385- '

_" 394 ..

5. Kear, B H. and B. J Piearcey: "Tensile and Creep Properties

2_

;_ of Single Crystals of the Nickel-Base Superalloy Mar-M 200," _

'_ Trans. AIME, 239, 1967, pp. 1209-1215. i

+

)

6. Leverant, G. R. and B. H. Kear: "The Mechanism of Creep in

Gamma Prime Precipitation-Hardened Nickel-Base Alloys at

Intermediate Temperatures," Metallurgical Transactions, i,

1970, pp. 491-498• _

: 7. Leverant, G. R., B. H. Kear, and J. M. Oblak: "Creep of Pre- !

cipitation - Hardened Nickel-Base Alloy Single Crystals at

+_ High Temperatures," Metallurgical Transactions, 4, 1973, pp.

: 355-362.

215

: 1

+ 1 + +.... ®
ii' ®

r

!

8. Boone, D. H., D. N. Duhl, and G. W. Goward: "Nickel-Base

Superalloy Resistant to Oxidation-Eroslon," U.S. Patent

3,754,902, August 28, 1973.

_. 9. Duhl, D. N. and X. Nguyen-Dinh: "Evaluation of Single-Crystal

Superalloy," presented at AIME Annual Meeting, Las Vegas,

February 24-28, 1980.

C

i0. Ver Snyder, F. L. and M. Gell: "New Direction in Alloy Design

for Gas Turbines," Fundamental Aspects of Structural Alloy

Design, Eds. R. I. Jaffee and B. A. Wilcox, Plenum Publishing

%. Corp., 1977, pp. 209-227.

ii. Gell, M., D. N. Duhl, and A. F. Giamei: "The Development of

Single Crystal Superalloy Turbine Blades, " in Superalloys _,

_ 1980- Proceedings of the Fourth International Symposium on

Superalloys, September 1980, Champion, Pennsylvania, pp. 205-

. 214.

::

12. Jackson, J. J., M. J. Donachle, R. J. Henrlcks, and M. Gell:

: "The Effects of Volume Percent oi" Fine Y' on Creep in DS Mar-M

t 200 + Hf," Metallurgical Transactions A, 8A, 1977, pp_ 1615-

1620 •

13. Strangman, T. E., G. S. Hoppln III, C. _,. Phipps, K. Harris,

and R. E. Schwer: "Development of Exothermlcally Cast SingleCrystal Mar-M 247 and Derlvatlve Alloys," in Superalloys

1980- Proceedings of the Fourth International Symposium on

Superalloys, Champion, Pennsylvania, September 1980, pp. 215-

224.

14. Strangman, T. E. and D. H. Boone: "Composition and Processing

_ Considerations for the Mechanical Behavior of CoatingSuperalloy Systems," in Proceedings of the Fourth Conference

on Gas Turbine Materials in a Marine Environment, Annapolis,

Maryland, June 1979, pp. 655-678.

b

. 216
®

!

I

15. Gell, M. and G. R. Leverant: "The Fatigue of the Nickel-Base

: Superalloy, Mar-M 200, In Single-Crystal and Columnar-Grained

Forms at Room Temperature," Trans. AIME, 242, 1968, pp. 1869-

_ 1879.

i

16. Gell, M. and G. R. Leverant: "The Characteristics of Stage

Fatigue Fracture in a High-Strength Nickel Alloy," ACTA MET.,

_ 16, 1968, pp. 553-561.

i:

:$

17. Leverant, G. R. and M. Gell: "The Influence of Temperature

and Cycllc Frequency on the Fatigue of Cube Oriented NickelBase Superalloy Slngle Crystals," Metallurgical Transactions,

A, 6A, 1975, pp. 367-371.

18. Puroshothaman, S. and J. K. Tien: "Slow Crystallographic Fatigue Crack Growth in a Nickel-Base Alloy," Metallurgical

Transactions, 9A, 1978, pp. 351-355.

19. Sadananda, K. and P. Shahinian: "Analysis of Crystallographic _

Hlgh Temperature Fatigue Crack Growth in a Nickel Base Alloy,"

Metallurgical Transactions, 12A, 1981, pp.343-351.

0

20. Strangman, T. E., E. J. Felten, and R. S. Benden: "Refinement

Q of Promising Coating Compositions for Directlonally Cast

Eutectics," NASA CR-135103, October 1976.

I

217

QUSGIIO_ |g14-- 751-100/519

J