NUREG/CR-6824 

 ANL-02/37 

 

 

Materials Behavior in

HTGR Environments 

Argonne National Laboratory

U.S. Nuclear Regulatory Commission

Office of Nuclear Regulatory Research

Washington, DC 20555-0001
 

 NUREG/CR-6824 

 ANL-02/37

Materials Behavior in

HTGR Environments

Manuscript Completed: February 2003

Date Published: July 2003

Prepared by

K. Natesan, A. Purohit, S. W. Tam

Argonne National Laboratory

9700 South Cass Avenue

Argonne, IL 60439

C. A. Greene, NRC Project Manager

Prepared for

Division of Engineering Technology 

Office of Nuclear Regulatory Research

U.S. Nuclear Regulatory Commission

Washington, DC 20555-0001

NRC Job Code Y6537 
ii 
iii 

Materials Behavior in HTGR Environments

by

K. Natesan, A. Purohit, and S. W. Tam

Abstract

The report reviews and evaluates data currently available on materials performance and

long-term behavior in High-Temperature Gas-Cooled Reactors (HTGRs), including the Gas

Turbine-Modular Helium Reactor (GT-MHR) and the Pebble Bed Modular Reactor (PBMR). As

part of the evaluation, the report identifies the materials that have been used in HTGRs, their

functions, and the environmental conditions under which they need to perform adequately.

The report identifies the materials used for structural applications (such as pressure vessel and

reactor primary circuit components, including internals) and for the power conversion system,

with emphasis on gas-turbine-based HTGRs. The report examines the chemistry of the

primary-circuit helium environment from the standpoint of type and concentration of

impurities. A framework is presented for analyzing gas-solid interactions in single and

bioxidant conditions, and the data on materials scaling behavior are compared from several

research programs. This report discusses the available data on the mechanical properties

(such as creep, fatigue, creep fatigue, and crack growth) of candidate materials in a helium

environment. Available information on the carburization, decarburization, oxidation, and aging

of structural materials and their effects on mechanical properties is summarized. From the

materials standpoint, the report also discusses the role of particulate-laden gases in corrosionerosion and the possible effects on materials from off-normal situations in the operation of the

reactors. The report makes several recommendations to supplement the data on

environmental effects on materials, on mechanical properties of structural and power

conversion materials, and on alternative candidate alloys for application in HTGRs.
iv 
v 

Contents

Abstract.................................................................................................................................... iii

Executive Summary................................................................................................................. xi

Acknowledgments .................................................................................................................... xv

1 Introduction .................................................................................................................... 1

1.1 Gas Turbine Modular Helium Reactor ............................................................... 3

1.2 Pebble Bed Modular Helium Reactor.................................................................. 7

2 Metallic Components ...................................................................................................... 9

2.1 Structural Materials............................................................................................ 9

2.1.1 2.25Cr-1Mo Ferritic Steel ..................................................................... 9

2.1.2 Modified 9Cr-1Mo Ferritic Steel............................................................ 10

2.1.3 Alloy 800H............................................................................................. 17

2.1.4 Alloy 617, Hastelloy X, and Other Alloys ............................................. 19

2.2 Power Conversion Materials ............................................................................... 24

2.2.1 Turbine Blade Alloys............................................................................. 25

2.2.2 Turbine Disk Alloys............................................................................... 25

3 Helium Coolant Effects ................................................................................................... 27

3.1 Coolant Chemistry .............................................................................................. 27

3.1.1 Equilibrium versus Nonequilibrium Gas Chemistry............................ 28

3.1.2 Effect of System Pressure ..................................................................... 29

3.2 Effect of Impurities on Corrosion ....................................................................... 29

3.2.1 Reaction in Single-Oxidant Environments........................................... 30

3.2.2 Reaction in Bioxidant Environments.................................................... 32

3.2.3 Corrosion Performance Data ................................................................ 35

3.3 Effect of Impurities on Mechanical Properties ................................................... 41

3.3.1 Creep and Creep Rupture Properties .................................................. 41

3.3.2 Fatigue Properties................................................................................. 45

3.3.3 Other Properties.................................................................................... 49

4 Corrosion-Erosion........................................................................................................... 50

5 Design Methodology........................................................................................................ 52

6 Off-normal Situations ..................................................................................................... 53

6.1 Loss of Coolant.................................................................................................... 53

6.2 Air and Water Ingress ......................................................................................... 54

6.3 Particulate-Laden Gases..................................................................................... 55

7 Summary and Recommendations .................................................................................. 57

References ................................................................................................................................ 63
vi 

Figures

1.1 Schematic of gas turbine-modular helium reactor ................................................... 4

1.2 Schematic of pebble bed modular helium reactor (PBMR) ....................................... 8

2.1 Rupture stress in development of 9%Cr steels ......................................................... 11

2.2 Secondary creep rates for P91, P92, and E911......................................................... 12

2.3 Comparison of ultimate tensile strength of unexposed Grade 91 with

aged and service-exposed materials.......................................................................... 12

2.4 Comparison of creep curves for unexposed and service-exposed Grade 91

at 100 MPa................................................................................................................. 13

2.5 Creep curves for service-exposed Grade 91 at 100 MPa and several

temperatures.............................................................................................................. 13

2.6 Comparison of isostress plots of minimum creep rate of Grade 91 for

several material conditions at 100 MPa.................................................................... 14

2.7 Comparison of isostress plots of rupture life of Grade 91 for several

material conditions at 100 MPa................................................................................. 14

2.8 Monkman-Grant plot for unexposed, aged, and service-exposed Grade 91

steel ............................................................................................................................ 15

2.9 Stress versus Larson-Miller parameter for unexposed, aged, and serviceexposed Grade 91 steel.............................................................................................. 15

2.10 Cyclic stress-strain properties of hot-rolled and forged materials of

modified 9Cr-1Mo steel.............................................................................................. 16

2.11 Fatigue crack propagation rate as a function of crack length for modified

9Cr-1Mo steel............................................................................................................. 16

2.12 Allowable stress values Smt for primary membrane calculations for actual

service life and for normal-plus-upset conditions, as specified by ASME

Boiler and Pressure Vessel Code............................................................................... 17

2.13 Minimum creep rate versus stress at 538-760°C and correlations

developed by regression analysis............................................................................... 18

2.14 Low- and high-cycle fatigue behavior of Alloy 800H at 760°C.................................. 18

2.15 Tensile properties of Hastelloy X and Inconel 617 in both solutionannealed and aged conditions................................................................................... 20

2.16 Applied stress versus time-to-1% strain from creep of Alloy 617 at several

temperatures.............................................................................................................. 21
vii 

2.17 Applied stress versus time-to-rupture from creep of Alloy 617 at several

temperatures.............................................................................................................. 21

2.18 Applied stress versus secondary creep rate for Alloy 617 at several

temperatures.............................................................................................................. 21

2.19 Larson-Miller parameter versus time-to-1% strain for Alloy 617............................. 22

2.20 Larson-Miller parameter versus time-to-rupture for Alloy 617................................ 22

2.21 Comparison of fatigue behavior of Alloy 617 with several conventional

materials at 538°C..................................................................................................... 23

2.22 Applied stress versus rupture time for Hastelloy X at 760 and 871°C .................... 23

2.23 Applied stress versus minimum creep rate for Hastelloy X at 760 and

871°C.......................................................................................................................... 24

2.24 Low-cycle fatigue behavior of Hastelloy X at several temperatures in air................ 24

3.1 Thermodynamic stability of various oxides as a function of temperature ............... 32

3.2 Thermodynamic stability of various carbides as a function of temperature............ 32

3.3 Thermochemical diagram for Cr-C-O system at 982°C, indicating the

stability region of various phases.............................................................................. 34

3.4 Different surface layers observed on high temperature alloys exposed to

helium in high temperature reactor .......................................................................... 35

3.5 Equilibrium partial pressure for several oxides as a function of reciprocal

temperature, along with calculated oxygen partial pressures, for

environments A through D based on Reaction (12) ................................................. 37

3.6 A comparative plot of corrosion behavior for Alloys 800H, Nimonic 86,

and 617 from 2- and 50-atm tests ............................................................................ 40

3.7 Comparative plot of the environments used in different research

programs listed in Table 3.2...................................................................................... 40

3.8 Creep rupture data for Alloy 800H at 650 and 760°C in helium with

comparison to specifications based on air data in ASME Code ............................... 41

3.9 Creep curves for Alloy 617 at 950°C in a helium environment................................ 42

3.10 Heat-to-heat variation in the creep behavior of Alloy 617 in helium ....................... 42

3.11 Variation in creep rupture time as a function of applied stress for

Alloy 617 in air and helium at several temperatures ............................................... 43

3.12 Variation in creep rupture time as a function of applied stress for

Alloy 617 at 950°C in air and helium of low steam content..................................... 43
viii 

3.13 Time-to-1% strain and time-to-rupture as a function of applied stress for

Alloy 617 in air and in HTGR helium........................................................................ 44

3.14 Creep rupture ductility for Alloy 617 in air and helium........................................... 44

3.15 Examples of creep curves and cracking behavior for Alloy 617 tested in

air and in helium at 950°C........................................................................................ 45

3.16 Comparison of low-cycle fatigue behavior of Inconel 617 tested in air and

in helium at 704°C..................................................................................................... 46

3.17 Comparison of low-cycle fatigue behavior of Inconel 617 tested in air and

in helium at 871°C..................................................................................................... 47

3.18 Comparison of the low-cycle fatigue behavior of Alloy 617 in air and

helium with the design curves................................................................................... 47

3.19 Comparison of low-cycle fatigue behavior of solution-treated and

carburized Alloy 617 tested in air and in helium at 1000°C.................................... 47

3.20 Influence of hold time on fatigue life of Alloy 617 at 1000°C ................................... 48

3.21 Comparison of experimental creep-fatigue data with linear-damage

accumulation rule...................................................................................................... 48

3.22 Comparison of strain-controlled fatigue data generated at 871°C in air

and in helium for Hastelloy X in several conditions ................................................. 49

6.1 Fission-product barriers and retention properties of TRISO fuel elements

for a pebble-bed gas reactor ...................................................................................... 54

6.2 Coated-particle fuel temperature capability.............................................................. 54
ix 

Tables

1.1 Characteristics of selected steam cycle HTGRs ........................................................ 2

1.2 Characteristics of selected modular HTGR gas turbine plants ................................ 2

1.3 GT-MHR component materials.................................................................................. 5

2.1 Nominal chemical composition of candidate alloys and their application............... 10

2.2 Chemical composition of several grades of 9Cr-1Mo steels...................................... 11

3.1 Range of impurity levels expected in HTGR helium coolants................................... 30

3.2 Gas chemistries used in various investigations........................................................ 36

7.1 Summary of available thermomechanical data for alloys considered for

pressure vessel and other non-core, lower temperature components...................... 60

7.2 Summary of available thermomechanical data for various alloys

considered for piping and internals........................................................................... 61

7.3 Summary of available thermomechanical data for various alloys for power

conversion applications ............................................................................................. 62
x 
xi 

Executive Summary

The High Temperature Gas Reactor (HTGR) concept currently considered by the Member

States of the IAEA's International Working Group on Gas Cooled Reactors (IWGGCR) involves

helium as the coolant and a closed-cycle gas turbine for power generation with a core

outlet/gas turbine inlet temperature of 850°C and a net plant efficiency of 47%. In these

designs, the helium from the core is used directly to drive the turbines or indirectly to heat the

air or nitrogen, which drives the turbines. The system concepts include the Gas TurbineModular Helium Reactor (GT-MHR) developed by a consortium led by General Atomics in the

U.S., the Pebble Bed Modular Reactor (PBMR) designed by ESKOM of South Africa and British

Nuclear Fuels of U.K., and the High Temperature Engineering Test Reactor (HTTR) of Japan.

The objective of this report is to review and evaluate the available information on

performance and on long-term behavior of materials in environments that are typical of hightemperature helium-cooled reactors. As a part of this evaluation, the materials that have been

used in HTGRs, their functions, and the environmental conditions under which they have

operated are identified. The materials to be used in new HTGRs, the functions they will serve,

and the anticipated environmental conditions under which they are expected to operate are

likewise identified. Those HTGR conditions that may affect material property response, such as

the steady-state operating environment, component geometry, and off-normal conditions, are

discussed and evaluated.

The primary helium coolant in the gas-turbine-based HTGRs is expected to be at

temperatures in the range of 850-900°C, and the selected materials should have adequate

performance over long service life at temperatures in the range of 900-950°C. There are several

issues in materials selection for such an application, and reliable component performance

requires a significant technical database for the selection of materials and for assessment of

the long-term effects on the materials behavior in the service environment. This report reviews

the literature on HTGR materials properties and environmental effects on the behavior of

metallic materials. The key technical issues addressed in this report are:

• Baseline mechanical property data,

• Effects of helium coolant chemistry on materials degradation,

• Corrosion effects on mechanical properties of candidate materials,

• Fission product release and its effect on materials,

• Corrosion-erosion due to particulate-laden gas flow.

Section 2 of the report discusses the baseline mechanical property data for several

candidate materials that are being considered for application in HTGRs. The properties include

tensile strength, creep rupture, low-cycle fatigue, creep fatigue, and fracture toughness. The

information is used to assess the adequacy and applicability of data for the design and service

life of components in HTGR environments. The availability and adequacy of design codes,

rules, and procedures for various components are discussed in a companion report.

Among the materials, 2 1/4Cr-1Mo and modified 9Cr-1Mo ferritic steels are considered

for application in reactor pressure vessels. Fe-Cr-Ni alloys such as Alloy 800H and austenitic

stainless steels are considered for recuperators and reactor internals. Alloys such as 617,
xii 

Hastelloy X, and Hastelloy XR are considered for components that will be exposed to helium

coolant at temperatures up to 900°C. Alloys such as 713LC and Mo-TZM are considered for

the turbine blade. Alloys such as A286, 706, and 718, are examined for turbine disk

application.

Section 3 examines the influence of helium coolant on the chemical compatibility of

structural materials that are planned for use in HTGRs. Helium, because of its chemical

inertness and attractive thermal properties, is used as a primary coolant in HTGRs. However,

the primary coolant in an operating HTGR is expected to be contaminated by small amounts of

gaseous impurities such as H2, H2O, CH4, CO, CO2, and O2 from a variety of sources in the

reactor circuit.

Several conclusions can be drawn from this evaluation:

• Most materials research and development programs, in support of HTGRs, were

conducted in the 1960s to early 1980s. The thrust of these programs was to develop

a database on materials for application in steam-cycle and process-nuclear-heat

based HTGRs.

• Very little work was done during this period on materials with emphasis on direct

and/or indirect gas-turbine-based HTGRs.

• The primary materials for high temperature application in HTGRs that have been

studied in detail are Alloys 800H and 617 and Hastelloy X. Among them, Alloy 800H

is code certified for temperatures up to 760°C for use in nuclear systems. A draft

code case for Alloy 617 has been developed. A substantial database has been

developed for both Alloys 800H and 617, and a limited database exists for

Hastelloy X. Since the high temperature scaling in Hastelloy X has not been

adequate, a modified version, Hastelloy XR, has been developed in programs

conducted in Japan.

• Even though helium, by itself, is inert towards the materials, it is often contaminated

by small amounts of gaseous impurities such as H 2, H2O, CH4, CO, CO2, and O2 from

a variety of sources in the reactor circuit. The gas chemistry and the thermodynamic

activity for carbon and oxygen in the gas phase are difficult to ascertain because of

the nonequilibrium nature of the gas mixture. Furthermore, most studies on gas

chemistry simulations were performed close to atmospheric pressure, whereas the

system pressure in the reactor is 7 MPa.

• Structural alloys can be significantly corroded by the gaseous impurities in helium at

elevated temperatures. Past studies have shown that the corrosion of heat resistant

materials such as austenitic stainless steels, Alloy 800H, and Alloy 617 may involve

oxidation, carburization, and decarburization depending on the exposure

temperature, carbon activity in the gas phase, and the alloy composition.

Furthermore, the corrosion process is “dynamic” in that it is dictated by the exposure

time, gas chemistry variations, integrity of the corrosion product scales, and presence

of particulates in the gas phase. A thermodynamic framework that includes

evaluation of chemistry under nonequilibrium conditions is presented in this report.

The available corrosion information for several candidate alloys exposed to helium
xiii 

environments, is evaluated to address the scaling behavior and its relationship to gas

chemistry.

• Available data are reviewed on the mechanical properties of candidate alloys in a

helium environment with low levels of impurities. In most of these tests the

environment contained sufficiently high H2O (typical of steam-cycle-based HTGRs),

and the alloys predominantly developed oxide scales on the surface during

mechanical testing. The mechanical properties evaluated included creep, creep

rupture, low- and high-cycle fatigue, creep fatigue, and fracture toughness.

• In studies that were conducted in a helium environment with reduced levels of H2O,

Alloy 617 did not develop protective oxide scales, and the creep properties of the alloy

showed a decrease in life when compared with data generated in air. The aluminum

content in Alloy 617 led to internal oxidation rather than an external scale, especially

in environments with low oxygen partial pressure. It is believed that the impurity

levels in gas-turbine based systems will be predominantly decarburizing since oxygen

potential is expected to be much lower than that in steam-cycle based systems.

• Before materials performance can be assessed, it is essential to establish an

operational window in terms of impurity levels in helium that can be maintained and

controlled with certainty.

• Over the last 20 years, alloy manufacturers have developed several alloys with

improved resistance to oxidation and to carburization/decarburization at elevated

temperatures. These materials include Alloy 230, 602CA, microalloyed cast HP alloys,

and W-containing Ni-base superalloys. It is beneficial to examine the scaling and

corrosion performance of these alloys (in addition to Alloys 617, Hastelloy X, etc.) in

helium with low oxygen partial pressures for their viability in the HTGRs.

• An extensive database is needed on candidate alloys for application in turbines (both

for blades and disks), especially for direct-cycle systems.
xiv 
xv 

Acknowledgments

This work is sponsored by the Office of Nuclear Regulatory Research, U.S. Nuclear

Regulatory Commission, under Job Code Y6537; Senior Technical Advisor: J. Muscara;

Program Manager: C. A. Greene. Figures 2.3 to 2.9, 3.9, 3.13, 3.14a, 3.18, and 3.21 in the

report are reprinted with permission from the American Society for Mechanical Engineers, who

is the original publisher of these figures. Figures 2.16 to 2.20, 2.22, 2.23, 3.6, 3.10, 3.11,

3.14b, and 3.15 in the report are reprinted with permission from the American Nuclear Society,

who is the original publisher of these figures. Figure 6.2 in the report is reprinted with

permission from General Atomics Company, who is the original publisher of the figure.

The sources for various figures and their authors are identified in the report along with the

figures.
1

1 Introduction

Development of the High Temperature Gas Reactor (HTGR) began in the 1950s, with

significant advances in system design and fuel concepts during the 1960s and 1970s (IAEA

1990). The designs may be divided into two main groups, namely, steam-cycle HTGRs and

modular-gas-turbine HTGRs. However, in all HTGR designs helium serves as the primary

coolant. Table 1.1 lists the characteristics of selected steam-cycle HTGRs. In the steam-cycle

HTGRs, helium from the primary circuit is used to produce steam, which in turn, drives the

turbines to generate power. The early HTGR designs included Dragon, Arbeitsgemeinshaft

Versuchsreaktor (AVR), and Peach Bottom, which utilized steel primary system vessels. The

Dragon reactor incorporated graphite fuel elements containing particles coated by highenriched-uranium carbide (Moore 1982). The core exit and inlet helium temperatures were 750

and 350°C, respectively.

In Germany, the 15 MWe AVR, a pebble-bed-type HTGR, began operation in 1967. The

AVR had a steel containment vessel and used particle-fueled, graphite spheres 6 cm in

diameter. The AVR operated with a core outlet temperature as high as 950°C (Kroger 1988).

The Peach Bottom HTGR (40 MWe) in the U.S. operated without fuel failures for 897 equivalent

full power days with fuel particles that were coated with isotropic pyrolytic carbon (Williams et

al. 1994).

The subsequent HTGR plants, Fort St. Vrain (FSV) and the thorium high temperature

reactor (THTR-300), had the primary systems enclosed in prestressed concrete reactor vessels

(PCRVs). The FSV reactor was designed with a core of hexagonal, graphite-block fuel elements

and reflectors with fuel in the form of ceramic-coated (TRISO) particles, once-through steam

generator modules producing 538°C superheated steam, and steam-turbine-driven axial

helium circulators (Moore 1982; Williams et al. 1994). The THTR-300 reactor, sponsored by

Germany and Norway, was also based on the pebble bed concept. The plant was completed in

1984, and the reactor demonstrated the desired safety characteristics, control response of the

pebble bed, and good fission product retention of the fuel elements before it was shut down

permanently in 1991. The designs discussed above are all based on steam plants for power

generation.

The advanced HTGR concept currently considered by the Member States of the IAEA's

International Working Group on Gas Cooled Reactors (IWGGCR) involves a closed-cycle gas

turbine for power generation with a core outlet/gas turbine inlet temperature of 850°C and a

net plant efficiency of 47% (Cleveland et al. 1997). Table 1.2 lists the characteristics of various

gas-turbine based HTGRs. In these designs, the helium from the core is used directly to drive

the turbines or indirectly to heat the air or nitrogen, which drives the turbines. The system

concepts include the Gas Turbine-Modular Helium Reactor (GT-MHR) pursued by a consortium

led by General Atomics in the U.S., the Pebble Bed Modular Reactor (PBMR) designed by

ESKOM of South Africa and British Nuclear Fuels of U.K., and the High Temperature

Engineering Test Reactor (HTTR) of Japan.
2

The main focus of this report is to evaluate the materials performance and long-term

behavior in environments typical of gas-turbine based HTGRs. As a part of this evaluation, the

materials that have been used in HTGRs, their functions, and the environmental conditions

under which they have operated are identified. The materials to be used in new HTGRs, the

Table 1.1 Characteristics of selected steam cycle HTGRs

(Boyer et al. 1974, Brey 2000, Nickel 1989)

AVR

Peach

Bottom

Ft. St.

Vrain Fulton THTR-300 HTR-500 VGM-400

HTRModule MHTGR

Country of

Origin

Germany U.S. U.S. U.S. Germany Germany Russia Germany U.S.

Thermal

Power,

MWt

46 115 842 2,979 750 1,390 1,060 200 350

Net.

Electric

Power,

MWe

13 40 330 1,160 300 550 Co-Gen. 80 139

Core

Outlet

Temp., °C

950 725 775 741 750 700 950 700 686

Helium

Pressure,

MPa

1.1 2.25 4.8 5.0 3.9 5.5 5.0 6.6 6.4

Steam

Temp., °C

505 538 538/538 513 530/530 530 535 530 538

Reactor

Type

Pebble Sleeve Block Block Pebble Pebble Pebble Pebble Block

Vessel

Material

Steel Steel PCRV* PCRV PCRV PCRV PCRV Steel Steel

Date of

Operation

1966 1967 1979 to

1989

No

operatio

n

1985 No

operatio

n

No

Data

No

operation

No

operatio

n

*Prestressed concrete reactor vessel.

Table 1.2 Characteristics of selected modular HTGR gas turbine plants* (Brey 2000)

GT-MHR PBMR MHTGR-IGT ACACIA GTHTR-300

HTR-GT

600MW MPBR

Country of Origin U.S./Russia S. Africa China Netherlands Japan Japan U.S.

Thermal Power, MWt 600 265 200 40 600 600 250

Net. Electric Power,

MWe

278 116 ~96 Co-Gen. 273 287 112

Core Outlet Temp., °C 850 900 900 800 850 850 850

Helium Pressure, MPa 7.15 7.0 6.0 2.3 6.8 6.0 7.9

Cycle Type

(Secondary Coolant)

Direct Direct Indirect

(nitrogen gas)

Direct Direct Direct Indirect

(air)

Core Type Block Pebble Pebble Pebble Block Pin/Block Pebble

Vessel Material Mod 9Cr1Mo SA 508 SA 516-70 Steel SA 508 Mod 9Cr1Mo Steel
3

*The systems were either planned or designed, but none was constructed.

functions they will serve, and the anticipated environmental conditions under which they are

expected to operate are likewise identified. Those HTGR conditions that may affect material

property response, such as the steady-state operating environment, component geometry, and

off-normal conditions, are discussed and evaluated.

1.1 Gas Turbine Modular Helium Reactor

The GT-MHR is an advanced gas-cooled reactor under development in a joint

United States-Russian Federation program. The GT-MHR is designed with the potential for

substantial improvements in the areas of safety, thermal efficiency, and environmental

advantage when compared with designs developed in the 1980s. The GT-MHR module, shown

in Figure 1.1, couples a gas-cooled modular helium reactor (MHR), contained in one vessel,

with a high-efficiency Brayton cycle gas turbine (GT) contained in a second vessel. The reactor

and power conversion vessels are interconnected with a short cross-vessel, and both the

vessels are located in a below-grade concrete silo.

Key features of the gas-cooled MHR are the helium coolant, graphite moderator, and

refractory-coated particle fuel. The helium coolant is relatively inert, the graphite moderator

has high strength and stability to high temperatures, and the refractory-coated fuel particle

retains the fission products to high temperatures. The helium coolant is heated in the reactor

core by flowing downward through the coolant channels in graphite fuel elements and then

through the cross-vessel to the power conversion system. The power conversion system

contains a gas turbine, an electric generator, and gas compressors on a common, vertically

oriented shaft supported by magnetic bearings. The power conversion system also includes the

recuperator, precooler, and intercooler heat exchangers. Table 1.3 lists the component

materials for GT-MHR plant and their operating environments.
4

Figure 1.1 Schematic of gas turbine-modular helium reactor (GT-MHR)
Table 1.3 GT-MHR component materials

Components Environment Material Comments References Code Status

Normal operating

temperature 290°C

Pressure ~7.07 MPa

SA 508 GR3,

Cl1

SA 508 Class 1

SA 533 Grade

B, Class 1

Insulated vessel design or

vessel cooled by helium

returning from turbine

EPRI 2002

Shenoy and Betts 1988;

Muto et al. 2000;

Buckthorpe et al. 2002

ASME Code Case

N-499-1

Reactor Pressure

Vessel

Cross-Vessel

Normal operating

temperature ~ 500°C

9Cr-1Mo-V

(also called

modified

9Cr1Mo)

(10X9MFB in

Russia)

This material is not yet

authorized as the Class 1

vessel material.

Mod 9Cr1Mo has a big

advantage over 2.25Cr1Mo

above 450°C; limited data for

as-received and post weld

heat-treated base material

and weldments; limited

fabrication experience. Tests

scheduled to begin mid 2002.

Muto et al. 2000

IAEA 2001

Buckthorpe et al. 2002

Kiryushin et al. 1997

None available

Reactor System

Metallic Components

- core support

 structure

- reactor internals

Core inlet/outlet

temperatures 491/850°C

2-1/4 Cr-1 Mo

Type 316 SS

Alloy 800H

Alloy 617

Hastelloy X

The support structure

includes the upper support

plate, core barrel shell, gas

duct shell, bottom collector

header shroud, hot duct

nozzle, and the shutdown

cooling system header. The

reactor internals include top

collector plenum and its

shroud, hot duct assembly,

control rod canister, plenum

element casing, and core

lateral restraint.

LaBar 2002;

Shenoy and Betts 1988

EPRI 2002

IAEA 2001

Natesan et al. 2003

ASME Code Case

N-201-4

(Maximum allowable

temperatures of

816°C or lower –

 See Table 5)

ASME Draft Code

Case for Alloy 617

(Maximum allowable

temperature of

982°C)

Shutdown Cooling

System (SCS)

- conical shell

- bottom

- gas circulator

850°C (conical shell)

490°C (bottom and gas

circulator)

Alloy

CrNi55MoWZr

(conical shell)

Type 321 SS

(gas circulator

blades)

EPRI 2002 None available

Reactor Core Cooling

System (RCCS)

Component

490°C Type 321 SS EPRI 2002 None available

5
Table 1.3 GT-MHR Component Materials (continued)

Components Environment Material Comments References Code Status

Hot Gas Pipe 850°C Alloy 617 ASME Draft Code

Case for Alloy 617

(Maximum allowable

temperature of

982°C)

Power Conversion

System (PCS) Vessel

150°C 15X2NMFA (in

Russia)

15Cr2NiMoVN

Kiryushin et al. 1997

EPRI 2002

ASME Section III,

Subsection NB

Recuperator (including

heat exchange surface)

Helium parameters for

low/high pressure side:

- Inlet 507.5°C, 2.63 MPa

 /108.1°C, 7.22 MPa

- Outlet 127.3°C, 2.60 Mpa

 /487.5°C, 7.15 MPa

Austenitic

stainless steel

08Cr16Ni11Mo3

IAEA 2001 ASME Section III,

Subsection NH

(Maximum allowable

temperatures of

816°C or lower – See

Table 5)

Precooler and

Intercooler

150°C (He)

60°C (air)

Type 321 SS EPRI 2002 ASME Section III,

Subsection NB

6
7

1.2 Pebble Bed Modular Helium Reactor

The PBMR is a helium-cooled, graphite-moderated, high-temperature reactor, a

schematic of which is shown in Figure 1.2. The PBMR design is aimed at achieving a plant

that has no physical process, however unlikely, that could cause a radiation-induced hazard

outside the site boundary (IAEA, 2001). This is principally achieved in the PBMR by

demonstrating that the integrated heat loss from the reactor vessel exceeds the decay heat

production in the post-accident condition, and that the peak temperature reached in the core

during the transient is below the demonstrated fuel degradation point and far below the

temperature at which the physical structure is affected.

The metallic parts in the core structure include the core barrel with its bottom support

structure, support flange at the lower section and stiffening ring at the top, helium flow

channels, pipes and components of the fueling system, fueling tube, and lower discharge tube.

The power conversion unit converts the heat from the reactor into electrical energy via a

Brayton cycle. The turbine inlet temperature is 900°C. A thermal power level of 265 MW is

calculated to provide a gross electrical generation of 117 MWe. Helium, at a pressure of

7.0 MPa and a temperature of 900°C, leaves the pebble bed reactor and flows through the highand low-pressure turbines, providing energy to drive their compressors. The helium, at 721°C

and 4.4 MPa, then enters the power turbine, providing energy to drive the electrical generator.

The helium at 554°C then enters the recuperator, which provides heat exchange to the helium

returning to the core. After leaving the recuperator at 140°C, the helium passes through the

precooler and exits the cooler at 27°C and 2.6 MPa. Before entering the high pressure side of

the recuperator, the helium is then compressed by passing through the low-pressure

compressor, intercooler, and then the high-pressure compressor. The recuperator heats the

helium from 104°C to 536°C at a pressure of 7.0 MPa before it enters the top of the reactor.
8

Figure 1.2 Schematic of pebble bed modular helium reactor (PBMR)
9

2 Metallic Components

The primary helium coolant in the gas-turbine based HTGRs is expected to be at

temperatures in the range of 850-900°C, and the selected materials should have adequate

performance over long service life at temperatures in the range of 900-950°C. There are several

key issues in materials selection for such an application, and reliable component performance

requires a significant database for the selection of materials and for assessment of the longterm effects on the materials behavior in the service environment. This report reviews the

literature on HTGR materials properties and the environmental effects on the behavior of

metallic materials. The technical issues addressed in this report are:

• Baseline mechanical property data developed in an air environment,

• Effects of helium coolant chemistry on materials degradation,

• Helium purification and control,

• Corrosion effects on mechanical properties of candidate materials,

• Fission product release and its effect on materials,

• Corrosion-erosion due to particulate-laden gas flow, and

• Design methodology.

2.1 Structural Materials

In this section, baseline mechanical property data are presented for several candidate

materials that are being considered for application in HTGRs. The properties include tensile,

creep rupture, low-cycle fatigue, creep-fatigue, and fracture toughness. The information is

discussed to assess the adequacy and applicability of data for the design of components in

HTGR environments with long service life. The availability and adequacy of design codes,

rules, and procedures for various components are discussed in a companion report (Shah et al.

2003).

Among the materials, 2 1/4Cr-1Mo and modified 9Cr-1Mo ferritic steels are considered

for application in reactor pressure vessels. Fe-Cr-Ni alloys such as Alloy 800H and austenitic

stainless steels are considered for recuperators and reactor internals. Alloys such as 617,

Hastelloy X, and Hastelloy XR are considered for components that will be exposed to helium

coolant at temperatures up to 900°C. Alloys such as 713LC and Mo-TZM are considered for

turbine blades. Alloys such as A286, 706, and 718 are examined for turbine disk applications.

Table 2.1 lists the nominal chemical compositions of several candidate alloys along with their

potential applications in HTGR components.

2.1.1 2.25Cr-1Mo Ferritic Steel

A substantial database is available on the tensile, creep, fatigue, and creep fatigue

properties for 2.25Cr-1Mo ferritic steel. In addition, this database includes the effect of longterm aging on microstructural changes and mechanical properties at temperatures of interest

for pressure vessel application in HTGRs. Furthermore, the material has been ASME-code
10

Table 2.1 Nominal chemical composition (in wt.%) of candidate alloys and their application

Material C Cr Ni Co Mo Ti Al Fe Other Application

2.25Cr-1Mo 0.20 2.3 - - 1.0 - – Bala – Pressure vessel

9Cr-1Mo 0.08 8.6 - - 1.0 - - Bal V 0.2, Nb 0.07 Pressure vessel

HT 9 0.2 12 0.3 - 1 - - Bal V 0.25, W 0.5 Pressure vessel

316 SS 0.05 18 14 - 2.2 - - Bal – Recuperator

800H 0.08 20.1 31.7 - 0.3 0.4 0.4 Bal – Piping, internals

Hastelloy X 0.1 21.5 Bal 2.0 9.0 - - 18.5 W 0.6 Piping, internals

617 0.06 21.6 53.6 12.5 9.5 0.3 1.2 0.9 – Piping, internals

Hastelloy-XR 0.07 22 Bal 0.04 9.1 - - 18.2 W 0.5, Mn 0.9, Piping, internals

Nimonic-86 0.07 25 Bal - 10 0.1 - 1.7 – Piping, internals

Hastelloy S 0.02 15.5 Bal - 14.5 - 0.2 – Piping, internals

Manaurite 0.4 25 35 - - - - Bal Nb 1 Piping, internals

Inconel 519 0.3 24 24 - - - - Bal Nb 1 Piping, internals

713LC 0.07 12 Bal - 5 0.7 6 - Nb 2, Zr 0.1 Turbine blades

Mar-M 004 0.07 12 Bal - 4.5 - 6 - Hf 1.3, Nb 1.6, Ta 0.3 Turbine blades

M21 0.1 6 Bal - 2 - 6 - W 10.5, Nb 1.5, Zr 0.1 Turbine blades

738 0.11 16 Bal 9.0 2 - 3.4 - W 2.6, Ta 2, Nb 0.9 Turbine blades

Udimet 520 0.05 19 Bal 12.0 6 3.0 2 - W 1 Turbine blades

Mo-TZM 0.02 - - - Bal 0.5 - - Zr 0.8 Turbine blades

A286 0.05 15 26.0 - 1.3 2.5 0.2 Bal B 0.015 Turbine disks

706 0.03 16 42.0 - 0.5 1.0 0.2 Bal Nb 2.9 Turbine disks

718 0.1 19 Bal - 3 0.8 0.6 19 Nb 5.2 Turbine disks

Waspaloy 0.08 20 Bal 13.5 4.3 3.0 1.3 - Zr 0.06 Turbine disks

Udimet 720 0.03 18 Bal 14.5 3 5.0 2.5 - W 1.3 Turbine disks

MA 6000 0.05 15 Bal - 2 2.5 4.5 - W 4, Ta 2, Y2O3 1.1 Turbine disks

aIndicates balance.

certified for application in nuclear systems. This report will not include a discussion of this

well-established database.

2.1.2 Modified 9Cr-1Mo Ferritic Steel

The 9Cr-1Mo-V steel is a modified alloy system similar to conventional 9Cr-1Mo grade

ferritic steel. Modifications include additions of vanadium, niobium, and nitrogen, as well as

lower carbon content. This alloy is much more resistant to thermal fatigue than austenitic

stainless steels because of its lower thermal expansion coefficient (at least 30% lower) and

higher thermal conductivity. This alloy provides extremely good mechanical properties at

elevated temperature when produced and heat treated to form the proper microstructure. This

microstructure includes stable submicroscopic vanadium-rich M23C6 particles that contribute

to high-temperature strength and creep resistance (Brush 2001).

Table 2.2 presents the chemical composition of three additional grades of 9Cr-1Mo-V

steels along with Grade P9 steel. The steel designated P91 was developed at Oak Ridge

National Laboratory (Sikka et al. 1981). It shows a remarkable increase in stress rupture

strength, achieved by addition of 0.2 V, 0.06 Nb, and 0.05 N. The steel designated P92 was

developed at Nippon Steel, in which a further increase in stress rupture strength was obtained

by addition of 1.8% W and reduction of the Mo content from 1 to 0.5%. Similar steel,

designated E911, was developed in a European program. This steel contains 1% Mo and 1% W

and offers stress rupture strength similar to P92. This development of 9Cr-1Mo-V steels is

illustrated in Figure 2.1, which shows the 100,000-h stress rupture strength of Grades P9,

P91, P92, and E911 at 600 and 650°C.
11

Table 2.2 Chemical composition (in wt%) of several grades of 9Cr-1Mo steels

Element P9 P91 P92 E911

C Max. 0.15 0.10 0.124 0.105

Si 0.20-0.65 0.38 0.02 0.20

Mn 0.80-1.30 0.46 0.47 0.35

P Max. 0.030 0.020 0.011 0.007

S Max. 0.030 0.002 0.006 0.003

Cr 8.5-10.5 8.10 9.07 9.16

Mo 1.70-2.30 0.92 0.46 1.01

W -- -- 1.78 1.00

V 0.20-0.40 0.18 0.19 0.23

Nb 0.30-0.45 0.073 0.063 0.068

B -- -- 0.003 --

N -- 0.049 0.043 0.072

Ni Max. 0.30 0.33 0.06 0.07

Al -- 0.034 0.002 --

Heat treatment 1 h at 1050°C + 1 h at

750°C, air cool

2 h at 1070°C + 2 h at

775°C, air cool

1 h at 1050°C + 1 h at

750°C, air cool

100,000-h stress

rupture strength

at 600°C (MPaa)

35 94 115 110

aValues for P91 from Canonico (1994), for P92 from Wachter et al. (1995), and for E911 from Staubli et al. (1998).

Figure 2.1 Rupture stress in

development of 9%Cr steels

(Sikka et al. 1983).

Extensive studies (Sikka et al. 1981, Brinkman et al. 1990, Swindemann et al. 1998)

have been conducted on a nominally Fe-9Cr-1Mo class of ferritic steels to evaluate their tensile,

creep rupture, and low-cycle fatigue properties. Long-term aging effects on the mechanical

properties have been determined with this material (Swindeman et al. 2000). Figure 2.2 shows

the secondary creep rates of P91, P92, and E911 at 600 and 650°C plotted against the applied

stress. In the high stress region, the differences in the secondary creep rates of the three steels

are relatively small, whereas the differences between the steels become more pronounced in the

low stress region. The creep deformation characteristics may be described by the Norton

equation with two values of Norton’s stress exponent n. At high stresses, the value of n is
12

around 16, whereas at low stresses, it is 6. The change in n indicates a change in the creep

characteristics.

Figure 2.2 Secondary creep rates

for P91, P92, and E911 (Ennis

and Czyrska-Filemonowicz 2002).

Figure 2.3 Comparison of ultimate tensile strength of unexposed Grade 91 with

aged and service-exposed materials (Swindeman et al. 2000).

Figure 2.3 shows the ultimate tensile strength after long-time aging, which is compared

to a trend curve for unexposed material. Plates aged for 75,000 h were tested at the aging

temperature. For aging at 482°C (900°F), little effect was noticed. If anything, a slight increase

in strength was observed. The ultimate strength at 649°C (1200°F) was significantly below the

trend curve for typical material. A more detailed analysis of the influence of aging time and

temperature is reported elsewhere (Brinkman et al. 1990). The material exposed to service

conditions for 130,000 h was tested at several temperatures. It was reported that the strength

fell below the trend curve for typical as-tempered material. At a comparable strain rate, the

strength of the service-exposed material was similar to the aged material at 593 and 649°C

(1100 and 1200°F).
13

Creep curves for service-exposed and unexposed modified 9Cr-1Mo steel are shown in

Fig. 2.4 for tests conducted at 600°C (1112°F) and 100 MPa (14.5 ksi). The service-exposed

material exhibited a creep rate that was ten times greater than the unexposed steel.

The unexposed material remained in the primary creep stage to beyond 15,000 h. Creep

curves for the service-exposed Grade 91 tested at different temperatures are shown in Fig. 2.5.

For 100 MPa, the creep rates at 625 and 600°C (1157 and 1112°F) accelerate around 1%

strain, whereas the creep rates at 575 and 550°C (1067 and 1022°F) showed no acceleration

around 1% strain after 20,000 h. The minimum creep rates at 100 MPa (14.5 ksi) for

three aging conditions are compared to the unexposed behavior in Fig. 2.6. These data are

plotted against the reciprocal of absolute temperature, and the trend indicates activation

energy as well as relative creep strength. The steeper slopes indicate higher activation

Figure 2.4 Comparison of creep

curves for unexposed and serviceexposed Grade 91 at 100 MPa

(Swindeman et al. 2000).

Figure 2.5 Creep curves for serviceexposed Grade 91 at 100 MPa and

several temperatures (Swindeman et

al. 2000).
14

Figure 2.6 Comparison of isostress

plots of minimum creep rate of

Grade 91 for several material

conditions at 100 MPa (Swindeman et

al. 2000).

Figure 2.7 Comparison of

isostress plots of rupture life of

Grade 91 for several material

conditions at 100 MPa

(Swindeman et al. 2000).

energies. The lowest value for reciprocal temperature corresponds to the highest temperature

of testing (700°C). With decreasing test temperature (increasing reciprocal temperature), data

tended to diverge. The service-exposed material exhibited the highest creep rate, and the data

for material aged 25,000 h at 649°C mingled within the trend. Material aged 10,000 h at

649°C exhibited lower creep rates, whereas the unexposed material exhibited the lowest rates.

For any temperature, the creep rates for all conditions varied by less than a factor of ten.

The slopes of the lines drawn through the data indicated that the activation energy for creep of

the service-exposed material was lower than that for the unexposed material. A plot of rupture

life versus the reciprocal temperature for the 100 MPa isostress data is shown in Fig. 2.7. With

the exception of one short-time test result on material aged 25,000 h at 649°C, the aged and

unexposed materials exhibited the same temperature dependence. The unexposed material

had the longest life, and material aged 25,000 h at 649°C had the shortest. The serviceexposed material showed a different temperature dependence. The slope of the line drawn

through the data indicated lower activation energy for rupture than the unexposed material.
15

The rupture lives for different conditions varied by less than a factor of ten at any temperature

examined. The rupture life of the pressurized tube was included in the plot, and the point fell

on the trend for the uniaxial tests.

Although the temperature dependence of the rupture life of service exposed material

appeared to differ from that for the aged materials, the Monkman-Grant correlation for all

conditions was similar. Data for several heats of unexposed material are compared to the aged

and service-exposed materials in Fig. 2.8. All data follow the same trend. As is often the case,

the trend for the long-time data was not as well defined as shorter time data. Rupture data for

more than forty tests on aged and service-exposed materials are plotted in Fig. 2.9 as stress

versus the Larson Miller parameter. A parametric constant of 30 was selected for

the correlation. Included in the plot are rupture data for a large number of tests on unexposed

material. Results clearly show that the rupture lives for the aged and service-exposed

materials fell near the lower strength or shorter life side of the scatter band for the unexposed

material. At the higher values of the Larson Miller parameter, the strength of the severely aged

materials was reduced to about 80% of the average for the unexposed material.

Figure 2.8 Monkman-Grant plot

for unexposed, aged, and

service-exposed Grade 91 steel

(Swindeman et al. 2000).

Figure 2.9 Stress versus

Larson-Miller parameter for

unexposed, aged, and serviceexposed Grade 91 steel

(Swindeman et al. 2000).

Monkman-Grant, isostress, and Larson-Miller correlations were used to describe the

strength of the material with aging in service in the temperature range of interest for pressure

vessel applications (Figs. 2.8 and 2.9). Changes in strength were attributed to coarsening of

the metallurgical substructure of the tempered martensite in the material (Swindemann et al.

2000).
16

Low-cycle fatigue behavior of modified 9Cr-1Mo steel has been evaluated in several

studies (Hoffman 1982, Jones 1983, and Ebi and McEvily 1984). The cyclic stress-strain

curves at 50% of the number of cycles to failure are shown in Fig. 2.10 (Ebi and McEvily 1984).

The results indicate a higher cyclic strength for the hot-rolled material compared with that for

the hot-forged material. Also included in Fig. 2.10 are correlations that relate cyclic stress and

cyclic strain. Ebi and McEvily (1984) also measured fatigue crack propagation rate as a

function of crack length by measuring the fatigue striation spacing on the failed test

specimens. They concluded that the crack propagation rate was similar for both the hot-rolled

and hot-forged materials when tested in air (see Fig. 2.11). But they reported that the crack

growth rates were lower in vacuum than in air for the hot-rolled material.

9Cr-1Mo-V materials have been adopted in various forms in a number of ASTM/ASME

standards. Many are listed in the ASME Boiler and Pressure Vessel Code Section II, Part D,

allowable stress tables with a maximum temperature rating of 649°C (1200°F) for ASME Code

Section I and Section VIII, Division 1, and rating of 371°C (700°F) for Section III, Class 1, 2,

and 3.

Figure 2.10 Cyclic stress-strain

properties of hot-rolled and forged

materials of modified 9Cr-1Mo steel

(Ebi and McEvily 1984).

Figure 2.11 Fatigue crack

propagation rate as a function of

crack length for modified 9Cr-1Mo

steel (Ebi and McEvily 1984).
17

2.1.3 Alloy 800H

Alloy 800H consists of Fe-Cr-Ni with a minimum carbon content of 0.05 wt.% and a grain

size of ASTM No. 5 or coarser. This alloy must be solution annealed at 1093°C, has stable

austenitic structure, and is intended for use at temperatures >593°C. Both the high carbon

content and the high solution-anneal temperature promote a large stable grain size, which

provides better creep properties. This alloy is currently approved under ASME Code for

nuclear service up to 760°C.

The material is free of precipitates in the solution-annealed condition. After long times at

a service temperature of 550°C, ' precipitates tend to form. These precipitates can reduce the

creep ductility of the alloy (Betteridge et al.1978). To minimize the decrease in creep ductility,

the volume fraction of ' precipitates is controlled by specifying a concentration limit of

0.8 wt.% for the Al + Ti content in the alloy. The minimum creep rate for the alloy at elevated

temperatures is depicted by Norton’s law and the Arrhenius function, as follows (Blackburn

1972):

 

ε  ̇ m = constant( /E)n exp(−Q/RT) (1)

where , E, Q, R, and T are applied stress, elastic modulus, activation energy, gas constant,

and absolute temperature.

Figure 2.12 shows the allowable stress values Smt used for primary general membrane

calculations for actual service life and for normal-plus-upset conditions, as specified by ASME

boiler and pressure vessel code (ASME 1977). Figure 2.13 shows the minimum creep rate

Figure 2.12 Allowable stress values

Smt for primary membrane

calculations for actual service life and

for normal-plus-upset conditions, as

specified by ASME Boiler and

Pressure Vessel Code (ASME 1977).
18

Figure 2.13 Minimum creep rate versus stress at 538-760°C (1000-1400°F)

and correlations developed by regression analysis (Booker et al. 1978).

versus stress at 593 to 760°C, with comparison to results derived by regression analysis

(Booker et al. 1978). Fatigue behavior of Alloy 800H has been evaluated from room

temperature to 760°C (Soo and Chow 1976, Soo and Chow 1978, Chow et al. 1978).

Figure 2.14 shows low- and high-cycle fatigue data (typical of information available) for Alloy

800H at 760°C, plotted as log-log plots of elastic and total strain range versus life.

Figure 2.14 Low- and high-cycle fatigue

behavior of Alloy 800H at 760°C ( Soo and

Chow 1978). BNL, GE, BMI, and GA in the

figures refer to Brookhaven National

Laboratory, General Electric Company,

Battelle Memorial Institute, and General

Atomics Company, respectively.
19

2.1.4 Alloy 617, Hastelloy X, and Other Alloys

The alloys discussed in this subsection are selected for service at temperatures in excess

of 750°C, and the components fabricated from these alloys will be subjected to exposures at

elevated temperatures to helium coolant with potential impurities. This subsection will

address the elevated-temperature mechanical properties of candidate alloys, and the effects of

helium coolant will be addressed in Section 3.

Alloy 617 is an austenitic alloy based on nickel-chromium with further alloying additions

of cobalt, molybdenum, and aluminum. The alloy has exceptional creep strength at

temperatures above 870°C, enhanced by solution strengthening from the molybdenum and

cobalt additions. It has good cyclic oxidation and carburization resistance, imparted by the

chromium and aluminum additions. The alloy retains toughness after long-time exposure at

elevated temperatures; it does not form embrittling phases such as sigma, mu, chi, or Laves.

It has good weldability and lower thermal expansion than most austenitic stainless steels. As a

result, the alloy has been extensively studied for application at elevated temperatures in

HTGRs.

Tensile properties of Alloy 617 and Hastelloy X in solution-annealed and aged conditions

were evaluated by Strizak et al. (1982), and the results are shown in Fig. 2.15. Several studies

have been conducted to evaluate the creep and rupture properties of Alloy 617 at temperatures

>650°C (Schubert et al. 1984, Tanabe et al. 1984, Cook 1984, Schneider et al. 1984, Osthoff et

al. 1984). A database has been developed for the alloy that includes (1) the creep properties

such as time for 1% strain and time to rupture, (2) the correlation between secondary creep

rate and applied stress, and (3) correlations using the Larson Miller parameter to depict time

for 1% creep strain and rupture as functions of test time and temperature. Data have been

generated up to test times of 20,000 h and in multiple heats of the alloy. Figure 2.16 shows

the experimental data on 1% creep strain as a function of applied stress at temperatures

between 800 and 1000°C. Figure 2.17 shows similar data for rupture time at several

temperatures. Figure 2.18 shows a plot that relates the secondary creep rate to the applied

stress at several temperatures. Figures 2.19 and 2.20 show correlations using the Larson

Miller parameter to depict time for 1% creep strain and rupture as functions of test time and

temperature.

Low-cycle fatigue, creep-fatigue, and crack-growth rate tests on Alloy 617 have been

conducted by several researchers to establish the susceptibility of the alloy to cyclic loading

and also to evaluate the role of surface interactions of coolant on the fatigue crack initiation

and propagation (Strizak et al. 1982, Meurer et al. 1984, Hsu 1991). Figure 2.21 shows the

fatigue behavior of Alloy 617 with several conventional materials at 538°C (Strizak et al. 1982).

Hastelloy X is a nickel-base superalloy with good oxidation resistance at temperatures up

to 1200°C and moderately good strength properties at temperatures up to 870°C. It is

essentially a single-phase austenitic alloy, which is a solid solution strengthened by additions

of Cr, Mo, and W. Several studies have been conducted in the U.S. and in Japan to evaluate

the tensile, creep rupture, low-cycle fatigue, creep fatigue properties, and fatigue crack growth

of Hastelloy X, specifically for application in the helium environment of an HTGR (Lee 1984,

Kurata et al. 1984, Nakanishi and Kawakami 1984, Tsuji and Kondo 1984, Bruch et al. 1984,

Krompholz et al. 1984, Strizak et al. 1982).
20

Figure 2.15 Tensile properties of Hastelloy X and Inconel 617 in both solution-annealed and

aged conditions (Strizak et al. 1982).
21

Figure 2.16 Applied stress versus

time-to-1% strain from creep of

Alloy 617 at several temperatures

(Schubert et al. 1984).

Figure 2.17 Applied stress versus

time-to-rupture from creep of

Alloy 617 at several temperatures

(Schubert et al. 1984).

Figure 2.18 Applied stress versus

secondary creep rate for Alloy 617

at several temperatures (Schubert

et al. 1984).
22

Figure 2.19 Larson-Miller parameter

versus time-to-1% strain for Alloy 617

(Schubert et al. 1984).

Figure 2.20 Larson-Miller parameter

versus time-to-rupture for Alloy 617

(Schubert et al. 1984).
23

Figure 2.21 Comparison of fatigue behavior of Alloy 617 with several conventional materials at

538°C. Lines represent best fit of experimental data (Strizak et al. 1982).

Figure 2.22 shows the data on rupture time variation with applied stress for Hastelloy X

at 760 and 871°C. Data generated for times up to 20,000 h showed that the rupture ductility

for the alloy is fairly high in an air environment (Lee 1984). Figure 2.23 shows the variation in

minimum creep rate as a function of applied stress for the alloy at 760 and 871°C. Figure 2.24

shows the low-cycle fatigue behavior of Hastelloy X at several temperatures in air (Strizak et al.

1982). Lines in the figure represent the best fit of experimental data.

Apart from Alloy 617 and Hastelloy X, several additional superalloys have been examined,

albeit to a lesser extent, for elevated temperature application in HTGRs. These alloys include

Nimonic-86, Hastelloy S, Manaurite-36X, and Inconel 519; available data on the mechanical

properties for these alloys are given in several of the references cited above for Alloy 617 and

Hastelloy X.

Figure 2.22 Applied stress

versus rupture time for

Hastelloy X at 760 and 871°C

(Lee 1984).
24

Figure 2.23 Applied stress

versus minimum creep rate for

Hastelloy X at 760 and 871°C

(Lee 1984).

Figure 2.24 Low-cycle fatigue behavior of Hastelloy X at several temperatures in air.

Lines represent best fit of experimental data (Strizak et al. 1982).

2.2 Power Conversion Materials

The GT-MHR module consists of a nuclear reactor as the source of heat and a power

conversion system (PCS) (IAEA 2000, Courturier and Escaravage 2002). The PCS receives the

thermal energy from the reactor system and converts it to usable electrical energy, with an

overall efficiency of >45%. The PCS includes a turbomachine, precooler, intercooler, and

connecting pipelines. The turbomachine consists of a turbocompressor, electrical generator,

bearings, and seals. The turbocompressor includes the turbine and compressor sections.

The helium received by the PCS is expanded through the gas turbine that is coupled to

the electrical generator. From the turbine exhaust, the helium flows through the hot side of

the recuperator, transferring heat to helium returning to the reactor. The helium leaving the

hot side of the recuperator is cooled before passing through the low-pressure compressor,

intercooler, and high-pressure compressor. The helium then passes in the cold side of the

recuperator, where it is heated for return to the reactor system.
25

The entire PCS assembly is installed in a vertical orientation and operates at a rotational

speed of 3000 rpm. In the GT-MHR concept, a 12-stage turbine with high efficiency is

proposed. With turbine inlet temperature of 850°C, blade and disk cooling are not considered

as necessary in the design (Courturier and Escaravage 2002). It was reported that the

materials for the turbine blades and disks should ensure a design life of 60,000 h, and the

alloys selected for the blade and the disk should have minimum strength of 225 and 180 MPa,

respectively, at 850°C (OKBM 1997).

2.2.1 Turbine Blade Alloys

Two classes of alloys, namely, nickel-base superalloys and molybdenum-based alloys, are

being considered for blade application. The nickel-based superalloys include 713LC, M21, and

MAR-M 004 (Courturier and Escaravage 2000).

Alloy 713LC is a cast nickel-base precipitation-hardened alloy that combines good creep

rupture characteristics and excellent strength and ductility at room temperature. In addition,

the alloy contains neither cobalt nor tantalum; these elements have been excluded for use in

the primary circuit because of potential radioactive contamination. The alloy is reported to be

well suited for the HTGR turbine blades, except for the first row of blades where cooling would

be necessary to achieve the required lifetime (Kohler 1995). Alloy M21 is a low-chromium

nickel-base cast alloy that combines precipitation hardening (Mo, Al, Nb) and solid-solution

hardening (W). Alloy MAR-M 004 is a Hf-modified version of 713LC in which the Hf seems to

improve the fracture toughness of the material.

Among the molybdenum-base alloys, Mo-TZM is a prime candidate that has been studied

from the standpoint of fabrication, heat treatment, tensile, and creep rupture properties.

An alloy with a composition of Mo-0.5 wt.% Ti-0.8 wt.% Zr has been produced by vacuum arc

melting and by powder metallurgy (Jakobeit 1983).

2.2.2 Turbine Disk Alloys

The materials for the turbine disk application include Inconel 718, Udimet 720, MA 6000,

A286, Waspaloy, and Inconel 706. The alloy Inconel 718 derives its high temperature strength

by precipitation of gamma-prime [Ni3(Al,Ti)] and gamma-double-prime (Ni3Nb) phases. The

time-temperature-precipitation plot of the alloy shows that both phases precipitate in the

temperature range of 700 to 900°C and that the (Ni3Nb) precipitates first while the Ni3(Al,Ti)

precipitates over longer aging times (Brooks and Bridges 1988). An extensive database is

available on the mechanical properties of Inconel 718 (Huntington Alloys 1978, Blatherwick

and Cres 1966). The data show that at a turbine inlet temperature of 850°C, the alloy would

require cooling to lower the temperature of the disks to 650°C to achieve the desired lifetime.

Alloy A-286 is a wrought iron-nickel '-strengthened superalloy that contains 2.1 wt.% Ti

and 0.3 wt.% Al. The ' phase Ni 3(Ti,Al) in this alloy is metastable and is reported to transform

from the strengthening cubic ' to a weakening plate-like hexagonal eta phase on exposure at

temperatures >650°C (Ross and Sims 1987). This limits the use of this alloy to a maximum

temperature of 650°C. Alternatively, alloys containing cobalt (such as Udimet 700 and 720)

generally possess high strength at elevated temperatures but need development for use as

uncooled blades with the desired lifetime.
26

Alloy MA 6000, a nickel-base alloy produced by mechanical alloying, is an oxidedispersion-strengthened superalloy with high strength and microstructural stability. The alloy

combines solution strengthening by tungsten and molybdenum, precipitate strengthening by

coherent ' (Ni3Al), and dispersion strengthening by yttria. Precipitate hardening enhances

strength at intermediate temperatures, and the dispersion hardening improves the strength at

temperatures >950°C. Creep properties of the alloy, evaluated in several studies, show that the

alloy can be used in the uncooled condition in HTGR, but substantial additional effort is

needed for producing large ingots with homogeneous composition by hot isostatic pressing or

hot extrusion and for acquiring long-term mechanical property data (Howson et al. 1980, Been

1987).
27

3 Helium Coolant Effects

The focus of this Section is to examine the influence of helium coolant on the chemical

compatibility of structural materials that are planned for use in HTGRs. Helium, because of its

chemical inertness and attractive thermal properties, is used as a primary coolant in HTGRs.

However, the primary coolant in an operating HTGR is expected to be contaminated by small

amounts of gaseous impurities such as H2, H2O, CH4, CO, CO2, and O2 from a variety of

sources, such as reactions of ingressed water and oil with core graphite, and outgassing of

reactor materials. These impurities are projected to be in ppm levels in the helium coolant, but

the upper bound would strongly depend on the level of purification used for the helium supply

and the leak tightness of the reactor system. Corrosion of structural alloys by these gaseous

impurities at elevated temperatures can be significant. Past studies have shown that the

corrosion of heat resistant materials such as austenitic stainless steels, Alloy 800H and 617

may involve oxidation, carburization, and decarburization depending on the exposure

temperature, carbon activity in the gas phase, and the alloy composition. Further, the

corrosion process is dynamic in the sense that it is dictated by the exposure time, gas

chemistry variations, integrity of the corrosion product scales, and presence of particulates in

the gas phase.

3.1 Coolant Chemistry

The helium coolant in an operating HTGR makes a complete circuit from the graphite

core to the heat exchangers or gas turbines and back to the core in several seconds. It is

reported that the gas components in the coolant, via reaction with the graphite in the core and,

to a limited extent, with themselves, will reach a steady state under this dynamic flow

condition and may approach an equilibrium state with respect to the core (Johnson and Lai

1981). Equilibrium between the surfaces of metallic components and the gaseous impurities in

the primary coolant helium is not expected to occur under these very fast flow conditions.

Therefore, the carbon activity and oxygen partial pressure in the helium coolant, under such

nonequilibrium conditions, will be determined by individual reactions that predominate in the

gas mixture.

Although the gaseous impurities in a primary coolant environment may not be in

equilibrium with themselves or with surfaces of metallic components, driving forces will exist

for gas-metal interactions to occur. The extent to which these interactions will occur will be

kinetic controlled, dictated by time, temperature, alloy chemistry, and surface condition of the

alloy. In such nonequilibrium conditions, potentials for gas-metal corrosion reactions may be

determined through equilibrium thermodynamics by considering each individual chemical

reaction that is possible between the metal and individual gaseous impurities.

From the structural materials standpoint, we are interested in reactions that can affect

the corrosion loss and/or influence the mechanical integrity; reactions that can lead to

processes such as oxidation, carburization, and decarburization are of interest. Carburization

and decarburization processes are determined by the carbon activity in the gas mixture relative

to that in the exposed metal surface. Similarly, the oxidation process is determined by the

oxygen partial pressure in the environment relative to the stability of oxides of the constitutive

elements that are present on the exposed metal surface.
28

Carbon activity in the gas phase may be determined by any of the reactions listed below:

CO + H2 = C + H2O, (2)

2CO = C + CO2, (3)

CH4 = C + 2H2, (4)

CH4 + CO = 2C + H2O + H2, (5)

CO = C + 1/2 O2, (6)

The carbon activity in the gas phase for each of the five reactions may be calculated from

its equilibrium constant (K) via the partial pressure (p) of each component:

aC = K1 (pCO · pH2) / pH2O for Reaction (2), (7)

aC = K2 (pCO)

2/ pCO2

 for Reaction (3), (8)

aC = K3 pCH4 / (pH2)

2 for Reaction (4), (9)

aC = (K4)

1/2

 · (pCH4

 · pCO / pH2O · pH2)

1/2 for Reaction (5), (10)

aC = K5 pCO / (pO2)

1/2 for Reaction (6). (11)

In the absence of thermodynamic equilibrium between all the species in the gas phase, the

carbon activity in the environment will be established by any of the reactions listed above, and

the dominant reaction is probably the one that gives the highest carbon activity.

The oxygen partial pressure in the helium environment will be determined by the following

reactions, in addition to Reaction (6):

H2O = 1/2O2 + H2, (12)

CO2 = 1/2O2 + CO. (13)

The oxygen partial pressure established by these reactions can be calculated as follows:

pO2 = (K5 · pCO/ aC)

2 for Reaction (6), (14)

pO2 = (K6 · pH2O/ pH2)

2 for Reaction (12), (15)

pO2 = (K7 · pCO2/ pCO)

2 for Reaction (13). (16)

3.1.1 Equilibrium versus Nonequilibrium Gas Chemistry

When equilibrium prevails between various gas species in a multicomponent gas mixture,

the carbon activity and oxygen partial pressure in the gas environment are determined from

the minimization of free energy for the gas mixture as a whole at a given temperature. This
29

means that all the species in the gas mixture are in thermodynamic equilibrium, and that the

kinetics of reactions does not play a role in determining the values. Therefore, the calculated

carbon activity and oxygen partial pressure are unique for a given temperature and are

independent of the reactions given above. Under these conditions, the interactions of the gas

species with material surfaces can be analyzed using equilibrium thermodynamics on gasmetal interactions to establish the surface reaction products as a function of gas chemistry,

temperature, alloy chemistry, and exposure time.

On the contrary, if the gas species in the multicomponent gas mixture are not in

equilibrium, then one or more of the reactions discussed above will influence the actual gas

composition (at a given temperature). Furthermore, the composition can change with time

influenced by the kinetics of these reactions. Therefore, determining which of the reactions

dictates the carbon activity and oxygen partial pressure is difficult, if not impossible, and

understanding the variation of these parameters with time is even more complex. In an HTGR

environment, the gas phase composition is probably determined by the level of impurities in

the helium at the reactor inlet, which is modified by reactions with graphite in the core at

elevated temperatures. Therefore, one can establish an anticipated envelope for the gas

chemistry, based on a mass balance and reaction equilibrium at the core temperatures. The

gas chemistry thus calculated may be modified along the primary loop by the reactions

discussed above, but the variation may be small, since the flow rate of helium is high and the

residence time is relatively small.

3.1.2 Effect of System Pressure

System pressure can significantly affect on the calculations of carbon activity and oxygen

partial pressure, depending on which of the reactions (discussed above) predominantly controls

the process. In experiments that are conducted at a total pressure of 1 atm, the partial

pressures and concentrations of the species are essentially the same. On the contrary, in a

high-pressure environment, the partial pressure of a species is the product of concentration of

that species and the system pressure. As a result, for a given gas composition, the carbon

activity calculated using different reactions discussed above is affected differently by the

system pressure. For example, carbon activity calculated by Eqs. (7) and (8) increases in

proportion to an increase in system pressure. Carbon activity calculated by Eq. (9) decreases

with an increase in system pressure while that calculated by Eq. (10) is unaffected by system

pressure. Carbon activity calculated by Eq. (11) increases with the square root of system

pressure.

For a given oxygen concentration, the oxygen partial pressure is calculated by Eqs. (14),

(15), or (16). The values calculated by Eqs. (15) and (16) are unaffected by variations in system

pressure, but the values calculated by Eq. (14) increase directly with an increase in system

pressure. The effect of system pressure on gas-metal reactions is not adequately addressed in

the literature.

3.2 Effect of Impurities on Corrosion

The basis for the selection of impurity concentration range is the HTGR systems that are

designed using steam cycles for generating electricity or process heat. The thinking on the

anticipated levels of impurities went as follows (Graham et al. 1981).
30

The helium entering the core has a certain amount of oxygen and H2O. First, the hot

graphite core removes virtually all free oxygen to form CO, but the slower kinetics of the

reaction of the graphite core with H2O leads to only partial reaction, and hence the helium

contains traces of H2O together with H2 and CO produced by the reaction of some of the H2O

with the core. Any CO2 which is present from degassing of the graphite reflectors is also

largely converted to CO by reaction with the hot core. The level of H2 is increased relative to

that expected from the reaction of H2O with the graphite by proton inleakage from boilers or

process reaction, but some H2 also reacts with the cooler parts of the core via a radiolytic

reaction to produce CH4. Oil inleakage can also lead to CH4 production. However, the CH4

tends to be removed by thermal cracking in the hotter parts of the core. The residual amount

in the helium depends on the actual peak core temperature, higher temperatures leading to

lower CH4 levels. Therefore, the main impurities in the primary circuit helium were identified

as H2, H2O, CO, CH4, and N2.

The actual levels of impurities in a given system depend on the detailed design and

construction features such as the purification ratio, the type of gas circulators, and the choice

of graphite/carbon used. The ranges expected in steam- and gas-turbine-based HTGRs, under

normal operating conditions, are estimated as shown in Table 3.1 (Graham et al. 1981).

The levels of impurities in the helium coolant will be different for the steam-generating

and direct-cycle gas-turbine HTGR systems. One of the major differences lies in the oxygen

potential (or oxygen partial pressure) of the environment. In the case of the steam-cycle system

where continuous ingress of low levels of water and steam into the primary coolant is possible,

the environment is expected to have a relatively higher oxygen potential (i.e., high ratios of H2O

to H2). In the gas-turbine-based system, it is anticipated that the helium coolant will be much

“drier” (i.e., low ratios of H2O to H2) owing to the lack of potential sources of continuous water

and steam ingress. With the understanding that the impurity types and concentrations differ

between the steam-based and gas-turbine-based systems, we will next examine the available

information on materials performance in such environments.

3.2.1 Reaction in Single-Oxidant Environments

Considerable research on the causes, effects, and prevention of different types of

corrosion has been underway for many years. However, when one studies reaction kinetics

Table 3.1 Range of impurity levels (in Pa) expected in HTGR helium coolants

Range of partial pressure (Pa)

Impurity Steam cycle Gas turbine cycle

H2 20 50

H2O 1 0.05

CO 1 5

CH4 2 5

CO2 <0.1 <0.1

N2 1.5 1.5

O2 <<<0.1 <<<0.1
31

that involves only one reaction equilibrium, such as H2O/H2 or CO2/CO for oxidation or

CH4/H2 for carburization, the reaction potential of the participating species can be established

by the standard free energy of formation for the reaction. For an oxidation reaction of a pure

metal M, such as

M + 1/2 O2 = MO, (17)

the oxygen partial pressure (pO2) for the M/MO2 equilibrium is given by

pO2 at M/MO2 phase boundary = exp( G°/RT), (18)

where G° is the free energy change for the M/MO2 equilibrium at temperature T. In binary

and ternary alloys, the activities of the reactive element (aM) and the oxide (aMO2) should also

be considered. In such cases, Eq. (18) becomes

pO2 at M/MO2 phase boundary = aMO2/ aM exp ( G°/RT). (19)

In general, element activity in an alloy is given by

aM = M·XM,

(20)

where M and XM are the activity coefficient and mole fraction of M in the alloy, respectively.

If M is not known, ideal behavior is assumed, and M is assigned a value of unity.

Under similar single-oxidant conditions, the oxidation or carburization of a metal can be

established by the standard free energy of formation of the metal oxide or carbide. Plots of

standard free energy of formation versus temperature for compounds such as oxides and

carbides are given in the literature and serve as a guide in establishing the relative stability of

the corrosion-product layers with respect to both the partial pressure of the reactive species in

the exposure environment and the alloy composition.

Figures 3.1 and 3.2 show the thermodynamic stability of several oxide and carbide

phases that form by reaction between oxygen or carbon and the constituents of various

structural materials such as heat-resistant alloys, superalloys, and refractory metals. As a

result, in classical studies on materials exposed to single oxidants, the central theme is to

understand the kinetic aspects of continuous scale development and the mechanisms of scale

growth as a function of variables such as temperature, time, and reactant activity in the

exposure environment. An enormous literature exists on nucleation and growth of corrosionproduct scales for a wide range of materials exposed to several single-oxidant environments. In

addition, the effort has concentrated on modifying alloy compositions to (a) form slower growing

scales (for example, alumina, chromia, or silica scales instead of iron oxides, NiO, or CoO in

heat-resistant iron-, nickel-, and cobalt-base alloys), (b) reduce corrosion rates (smaller

parabolic rate constants), (c) improve adhesion of scales to substrates (by additions of reactive

elements such as Y, Ce, and La), especially under thermal cycling conditions, and (d) provide

surface enrichment of structural alloys with elements (such as Cr, Al, and Si on low- and

intermediate-chromium steels, and austenitic alloys) that form slow-growing oxide scales.
32

-140

-120

-100

-80

-60

-40

-20

400 600 800 1000 1200

Free Energy Change (Kcal/g atom O)

Temperature (°C)

Cr2O3

FeO

NiO

NbO

SiO2

MoO2

V9O

Ti3O2

ZrO2

Y2O3

Figure 3.1 Thermodynamic

stability of various oxides as a

function of temperature.

-50

-45

-40

-35

-30

-25

-20

-15

-10

400 600 800 1000 1200

Free Energy Change (Kcal/g atom O)

Temperature (°C)

Mo2C

Cr23C6

SiC

V2C

TiC

ZrC

Figure 3.2 Thermodynamic

stability of various carbides as

a function of temperature.

3.2.2 Reaction in Bioxidant Environments

When materials are exposed to multicomponent gas environments or mixed oxidants, the

oxidation and carburization processes mentioned above can occur simultaneously.

Furthermore, the second reactant may either modify or degrade the corrosion-product layers

that form by reaction of the first reactant and the substrate elements. In these situations,

reactions will occur not only at the gas/scale interface, but also within the scales that develop

on the surface and in the base metal beneath the reaction products. As a result, the

development of suitable and meaningful thermodynamic and kinetic frameworks for both

nucleation and growth of scale layers on materials exposed to mixed-oxidant atmospheres is

quite complex.

When studying reactions between metallic alloys and complex gas environments, the

chemical activities of several reactive species in the gas phase must be considered

simultaneously, and these activities are generally established by gas-phase equilibria,

especially at elevated temperatures. In binary and ternary gas mixtures, the chemical

potentials of the reactive species can be readily established, as a function of temperature, from
33

the room–temperature gas composition and free energy data for reactions. For a gas system

that consists of CO, CO2, H2, H2O, CH4, and N2, several gas equilibria have been used to

formulate a set of nonlinear algebraic equations (Natesan 1976, Chopra and Natesan 1977).

One can use iterative procedures to determine the gas compositions at elevated temperatures

that yield minimum free energy for the system and satisfy the conservation of different reactive

elements (i.e., C, H, O, and N) and the total pressure of the gas mixture. Results of such

analyses have been used to establish the elevated-temperature gas compositions, the partial

pressures of oxygen and nitrogen, and the carbon activity in the gas mixture. Analyses of

gas/metal interactions in mixed-gas atmospheres assume that equilibria among various

molecular gas species prevail, and chemical activities are calculated by using the above

approach or some similar method. However, under nonequilibrium conditions, the oxygen

partial pressure and carbon activity will be determined by any of the possible reactions, as was

discussed in Section 3.1.

We next address the thermodynamic aspects of metal corrosion in a bioxidant situation,

such as an oxygen/carbon environment. The following surface reactions are possible on a

divalent metal A:

A + 1/2 O2 = AO, (21)

A + C = AC, (22)

where AO and AC are the oxide and carbide reaction products, respectively. Under equilibrium

conditions, the partial pressures of oxygen and activity of carbon are defined by the relations

(aO)

eq

 = pO2

1/2 = exp ( GAO/RT), (23)

(aC)

eq

 = exp ( GAC/RT), (24)

where GAO and GAC are the standard free energies of formation of the oxide and carbide,

respectively, at temperature T. From Eqs. (23) and (24), one should be able to deduce the

conditions for oxidation or carburization; however, a further reaction must be considered,

namely

AC + 1/2 O2 = AO + C, (25)

with the equilibrium condition

(aC/aO)

eq

 = exp [( GAC – GAO)/RT] aAC/aAO. (26)

If we assume unit activity for the phases AC and AO, Eq. (26) is reduced to

(aC/aO)

eq

 = exp [( GAC – GAO)/RT]. (27)

Examination of Eqs. (23), (24), and (27) permits the identification of various situations

that limit the type of surface corrosion products that can be formed, as follows:

(i) if (aO)

gas 

> (aO)

eq 

and

 

(aC)

gas 

< (aC)

eq

, (28)
34

then AO is the only stable surface phase;

(ii) if (aO)

gas 

< (aO)

eq 

and

 

(aC)

gas 

> (aC)

eq

, (29)

then AC is the only stable surface phase; and

(iii) if (aO)

gas 

> (aO)

eq 

and

 

(aC)

gas 

> (aC)

eq

, (30)

then both AO and AC should be stable and form as surface products. However, reference to

Eq. (25) indicates that only one phase will form, depending on which of the following conditions

prevails:

(a) (aC/aO)

gas

 > (aC/aO)

eq

, (31)

(b) (aC/aO)

gas

 < (aC/aO)

eq

. (32)

Condition (a) will cause Eq. (25) to proceed to the left, and AC will be the stable phase

where the metal is in contact with the gas phase. In condition (b), AO will be the stable phase,

and Eq. (25) will proceed to the right.

A convenient way of representing the possible corrosion products as a function of gas

chemistry is to construct thermochemical diagrams that depict the stability ranges of various

condensed phases as functions of the thermodynamic activities of the two components of the

reactive gas. Figure 3.3 shows a thermochemical stability diagram for the Cr–C–O system

developed for a temperature of 982°C. In the construction of this diagram, the thermodynamic

activities of the metal and corrosion–product phases are assigned a value of unity. In

multicomponent alloys, the activities of constituent elements will be less than unity and should

be accounted for in the analysis. Further, in gas/solid reactions that involve multicomponent

alloys, a more complex corrosion product (i.e., more complex than a binary compound) can

result that would decrease the thermodynamic activities of the specific phase in the mixture.

Figure 3.3 Thermochemical

diagram for Cr-C-O system at

982°C, indicating the stability

region of various phases.
35

3.2.3 Corrosion Performance Data

A significant body of corrosion information has been developed on several candidate

alloys exposed to helium environments with a narrow range of impurity concentrations

(Brenner and Nilsen 1978, Graham et al. 1981, Johnson and Lai, 1981, Kondo 1981, McKee

and Frank 1981, Bates 1984, Cappellaere et al. 1984, Kondo 1984, Huchtemann 1989, and

Graham 1990). The selected gas compositions were based predominantly on steam-cycle

HTGRs, which were envisioned during the 1970s and 1980s. A listing of gas compositions

used in several of the investigations is given in Table 3.2. Even though several studies have

been conducted to evaluate the corrosion performance of structural materials in helium that

contained different concentrations (in a somewhat narrow range) of impurities such as H2,

H2O, CH4, CO, and CO2, none of the studies was conducted with a systematic variation in gas

composition to evaluate its influence on the scaling and its protective capacity at elevated

temperatures. Figure 3.4 shows examples of scale morphologies that have been observed on

Alloy 800H, Nimonic 86, and Alloy 617. There are variations in scaling depending on alloy

chemistry. Scaing can also be influenced by gas chemistry in the exposure environment and

temperature.

Johnson and Lai (1981) examined the carburization behavior of Alloys 800H and 617 and

Hastelloy X in helium containing various amounts of H2, H2O, CH4, CO, and CO2, in a

temperature range of 649 to 1000°C for exposure times up to 10,000 h. Four different helium

environments, identified as A, B, C, and D (see Table 3.2), were used in the experiments.

Environments A and B were characteristic of high-oxygen potential, whereas environments C

and D were characteristic of low-oxygen potential. Thermodynamic calculations indicated that

all four environments were reducing to iron oxides (and to thermodynamically less stable

oxides of nickel, molybdenum, cobalt, etc.) and oxidizing to chromium and to elements more

reactive than chromium (manganese, silicon, titanium, aluminum, etc.). Figure 3.5 shows the

equilibrium partial pressure for several oxides as a function of reciprocal temperature, along

with the calculated oxygen partial pressures, for environments A through D based on

Reaction (12). Several conclusions were reported based on the results from the study.

Figure 3.4 Different surface layers observed on high temperature alloys exposed to helium in

high temperature reactor (Schubert 1984)
36

Table 3.2 Gas chemistries used in various investigations

Gas Gas composition (in Pa)

Identification H2 CO CH4 CO2 H2O N2 Reference

A 150 45 5 <0.05 5 - Johnson and Lai 1981

B 20 10 2 0.5 5 - Johnson and Lai 1981

C 50 5 5 <0.05 <0.05 - Johnson and Lai 1981

D 50 5 5 <0.05 0.15 - Johnson and Lai 1981,

Brenner and Nilsen 1978

HTR 50 5 2 - 0.1 - Graham et al. 1981

Decarb. 50 0.5 0.5 - 0.1 - Graham 1990

Carb. 50 0.15 0.2 - 0.01 - Graham 1990

JAERI-B 20 10 0.5 0.2 0.1 <0.5 Kondo 1981

ERANS#2 30 10 0.4 0.1 0.3 <0.5 Kondo 1981

PNP 50 2 2 0.1 <0.2 0.5 Kondo 1981, Quadakkers and

Schuster 1984, Bates 1984

GE 40 4 2 0.02 0.2 0.6 McKee and Frank 1981

Lab I 150 45 5 - 5 - Cappellaere et al. 1984

AIDA 150 45 5 - 5 - Cappellaere et al. 1984

Lab II 50 1.5 2 - 0.15 0.5 Huchtemann 1989

ERANS: Engineering Research Association of Nuclear Steelmaking, Japan.

JAERI: Japanese Atomic Energy Research Institute.

HTR: High Temperature Reactor.

PNP: Prototype Nuclear Process.

GE: General Electric Company.

Lab I and Lab II: Environments in laboratory tests.

• Carburization was observed in all four environments. For all three alloys,

carburization kinetics in the low-oxygen-potential environments (C and D) were

significantly higher (approximately one order of magnitude) at elevated temperatures

than those in the high-oxygen-potential environments (A and B) for all three alloys.

The enhanced carburization was attributed to kinetic effects in low-oxygen-potential

environments rather than the variation in gas chemistry.

• The scale formed on the metal surface in the high-oxygen-potential environments

consisted of oxides of manganese and/or chromium, while the scale formed in the

low-oxygen-potential environments consisted of Mn- and/or Cr-rich oxides and

Cr-rich carbides.

• Controlling the oxygen potential of the service environment can be an effective means

of reducing carburization of austenitic alloys. The effect of increasing the oxygen

potential on the corrosion rates of HTGR core graphite must be considered in

evaluating the practical application of this approach.

Graham et al. (1981) examined Alloy 617 and Nimonic 86 in a helium environment that

was considered as a standard gas (see Table 3.2) for helium-cooled process heat systems by
37

Figure 3.5 Equilibrium partial pressure for several oxides as a function of

reciprocal temperature, along with calculated oxygen partial pressures, for

environments A through D based on Reaction (12). Data from Johnson and

Lai 1981.

exposing the alloys at temperatures between 800 and 1000°C for times up to 10,000 h. They

related the degree of depletion or production of impurities caused by the gas/metal reactions

that occur at the metal surface by a parameter relating the surface area of the specimen and

volumetric flow of helium. The study reported several conclusions, as follows:

• Both the alloys developed thin, adherent, chromium-rich oxide scales. These scales

gave excellent resistance to carburization, although scale variations occurred due to

the influence of minor alloying elements such as titanium, aluminum, and

manganese.
38

• The influence of variations in gas chemistry showed that if the CH4/H2O ratio is

unfavorable, enhanced carburization resulted, and if the CO/H2O ratio is

unfavorable, decarburization was noted.

In a further study, Graham (1990) evaluated the long-term performance of Alloys 800H

and 617 at temperatures between 400 and 1000°C from the standpoint of decarburization,

carbon deposition/carburization, internal oxidation of the substrate, and spalling of the

surface scales. Two different gas chemistries were used, identified as “Decarb” and “Carb” in

Table 3.2. Several conclusions were reported from this study:

• Below 475°C corrosion is minimal, and the high-chromium alloys developed thin

chromia scales. At temperatures up to 900°C, the high-chromium alloys exhibited

thin chromia scales along with internal oxidation of aluminum and precipitate-free

zones. Carburization effects are enhanced if the scale is porous or if it spalls. Above

900°C rapid decarburization or carburization occurred, depending on the gas

chemistry.

• The onset of rapid carburization or decarburization occurred above a critical

temperature that is dependent on the alloy composition and the CO partial pressure

in the environment. It was postulated that at this critical temperature CO becomes

relatively more stable than the metal oxide and metal carbide. Above the critical

temperature, the H2 in the environment reduces the chromia scale, and the released

H2O reacts with carbon in the alloy to form CO. Graham (1990) termed this

mechanism as “gas catalyzed” reaction, whereas Brenner (1982) called it a

“microclimate reaction.” Brenner concluded that a “low” CH4/H2O ratio (i.e., excess

H2O) stabilized the Cr-based surface oxide and led to decarburization, whereas a

“high” CH4/H2 ratio (i.e., excess CH4) led to Cr-based surface carbide and continuous

carburization.

Kondo (1981) evaluated the corrosion performance of Hastelloy X and XR in two gas

mixtures (see Table 3.2) and concluded that Hastelloy X suffered from intergranular attack and

internal oxidation. The modified alloy XR did not show appreciable susceptibility to such an

attack, and it formed a continuous oxide film up to 1000°C in both environments. The kinetics

of Cr depletion followed a parabolic rate law, based on exposures up to 6,000 h at 1000°C. The

enhanced resistance of Hastelloy XR was attributed to formation of a double layer that

consisted of an outer layer of MnCr2O4 and an inner layer of Cr2O3. Kondo also suggested that

it is essential to maintain the ratio of the reacting metal surface to the volumetric flow rate of

helium-containing impurities during experimentation, thereby minimizing the change(s) in gas

chemistry during long exposure periods.

Bates (1984) established the long-term corrosion behavior of Alloys 800H and 617,

Nimonic 86, and Hastelloy X for the primary circuit in the prototype nuclear process (PNP)

reactor. Exposures were conducted in the temperature range of 700 to 900°C for times up to

10,000 h. In PNP helium chemistry (see Table 3.2), Nimonic 86 exhibited resistance to

carburization at temperatures of 700 to 900°C, and this behavior was attributed to formation of

defect-free chromia scales. Alloys 800H and 617 and Hastelloy X also showed good

carburization resistance, but a slow increase in carbon level with a parabolic kinetics. This

behavior was attributed to interference (in the formation of chromia scales) from other oxideforming elements in the alloys. Complete resistance to spalling of the scales was reported in all
39

alloys, except Alloy 800H. The cause for severe spalling in Alloy 800H was attributed to the

presence of 0.5 wt.% Si in the alloy. The good carburization resistance observed in the alloys

was lost or severely reduced when water levels in the gas dropped below 0.05 Pa.

McKee and Frank (1981) examined the corrosion behavior of several experimental alloys

in a simulated helium environment (see their Table 7 for gas composition) to evaluate the role

of alloying additions such as Al, Ti, Si, Nb, and Y in a base composition of Ni-20 wt.% Cr.

Experiments were conducted for 1000 h at 750 and 850°C and for 1000 and 3000 h at 950°C.

The alloys exhibited considerable variation in corrosion behavior over these test conditions. In

general, carburization was more severe at 950°C than at lower temperatures. Alloys containing

Al were susceptible to internal oxidation at all three temperatures and to carburization at

950°C. Additions of Nb promoted formation of protective oxide scales at all three temperatures.

Alloys containing Si were resistant to carburization but the scales tended to spall. Alloy

containing 1 wt.% Y exhibited massive precipitation of Cr carbides and poor corrosion

performance at 950°C.

Cappellaere et al. (1984) examined the influence of helium pressure on the corrosion of

ferritic and austenitic materials by conducting tests at 2 atm in a circuit without helium

recirculation and at 50 atm in the AIDA loop. The gas chemistries in both these experiments

were the same (Table 3.2). Experiments were conducted with HT 9 ferritic steel, Types 304 and

316 austenitic stainless steels, Alloys 800H and 617, and Hastelloy X. Specimens were

exposed at temperatures between 550 and 870°C for time periods up to 10,000 h at 2 atm and

for 2,500 h at 50 atm. Several conclusions were drawn from the study:

• Up to 650°C, HT 9 ferritic steel and Types 304 and 316 austenitic steels, and

Alloy 800H showed excellent corrosion performance.

• At 750 and 870°C, Hastelloy X exhibited better resistance to oxidation than

Alloys 800H and 617. For all three alloys, the oxidation rates decreased with time, but

the rates were 3 times higher for Alloys 800H and 617 when compared with

Hastelloy X.

• In general, the oxidation rates at 50 atm were higher than those at 2 atm. At 870°C,

the materials were more damaged after 2,500-h exposure in 50-atm tests, compared

with 10,000-h exposure in 2-atm tests. Figure 3.6 shows a comparative plot of the

corrosion behavior for the three alloys from 2- and 50-atm tests.

Huchtemann (1989) reported that in a helium environment with low oxygen pressure (see

Table 3.2 for gas chemistry used), the Alloy 617 surface underwent an enrichment of titanium

and had no manganese chromium spinel. In addition, an enhanced decarburization and

internal oxidation of aluminum occurred.

Figure 3.7 is a comparative plot of the environments used in different research programs

listed in Table 3.2. The carbon activity in this plot is based on Eq. (4), whereas the pO2 is

based on Eq. (12). Consistent with observations reported in several studies, a tentative line

can be drawn depicting the transition from the carburizing to oxidizing environment for high

chromium alloys. Even among the oxidizing environments, the alloy may undergo

decarburization, if the protective scaling process is kinetically slow. The results presented here

clearly show that significant additional tests are needed in well-controlled environments to
40

establish the mode of interaction for materials and its dependence on alloy and gas chemistry,

exposure time, and temperature.

Figure 3.6 A comparative plot of corrosion behavior for Alloys 800H,

Nimonic 86, and 617 from 2- and 50-atm tests (Cappellaere et al. 1984).

10-5

10-4

10-3

10-2

10-1

10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1

Parameter proportional to aC

Parameter proportional to pO2

A, LabI

B

C D

HTR

Decarb

Carb

JAERI, GE

ERANS

PNP LabII

AIDA

Figure 3.7 Comparative plot of the environments used in different

research programs listed in Table 3.2. The carbon activity in this plot is

based on Eq. (4), whereas the pO2 is based on Eq. (12).
41

3.3 Effect of Impurities on Mechanical Properties

Several studies have been conducted to evaluate the mechanical properties of candidate

alloys for structural applications and for power conversion systems. The mechanical properties

evaluated in these studies include creep, creep rupture, low- and high-cycle fatigue, creep

fatigue, and fracture toughness. Mechanical tests were conducted in a helium environment

that contained impurities such as H2, H2O, CH4, CO, and CO2. The impurity levels were

selected to simulate HTGR environments that contained sufficiently high H2O (typical of steamcycle-based HTGRs) so that the alloys predominantly developed oxide scales on the surface

during mechanical testing.

3.3.1 Creep and Creep Rupture Properties

Several creep studies have been conducted on Alloys 800H, 617, Hastelloy X,

Hastelloy XR, Nimonic 86, Manaurite 36X, and IN 519 in a “typical” reactor-helium chemistry

that was estimated for steam cycle and process heat applications (Schubert 1984, Schubert et

al. 1983, Schubert et al. 1984, Lee 1984, Kurata et al. 1984, Tanabe at al. 1984, Nakanishi

and Kawakami 1984, Cook 1984, Ennis et al. 1984). Most of the studies have emphasized

testing of Alloys 800H, 617, and Hastelloy X.

The mechanical properties, influence of the helium environment on the corrosion and

mechanical performance, and design aspects of Alloy 800 are discussed in detail elsewhere

(Betteridge, 1978) and are not elaborated further in this report. This report concluded that the

influence of helium environment (with a chemistry typical of steam-based HTGRs) on the creep

properties of the alloy is minimal at temperatures up to 760°C (see Figure 3.8), but that more

work is needed to fully characterize the behavior at higher temperatures and longer times.

The creep curves of Alloy 617, especially at elevated temperatures of interest in HTGRs,

show a monotonically increasing creep rate from the start of the test and virtually no secondary

creep region (see Figure 3.9). This behavior is different from that of austenitic stainless steel,

which distinctly shows primary, secondary, and tertiary creep regimes.

Figure 3.8 Creep rupture data for Alloy 800H at 650 and 760°C in helium with comparison to

specifications based on air data in ASME Code (Roberts 1978).
42

Figure 3.9 Creep curves for Alloy 617 at

950°C in a helium environment

(Schubert et al. 1983).

Figure 3.10 Heat-to-heat variation in

the creep behavior of Alloy 617 in

helium (Ennis et al. 1984).

Therefore, the creep data for Alloy 617 rarely include strain and time for onset-of-tertiary creep.

At elevated temperatures, the creep properties of interest are the minimum stress for 1% creep

strain and the rupture strength (Schubert 1983). In addition, the creep behavior of the alloy

seems sensitive to factors other than composition, as evidenced by creep test data generated on

three heats of the alloy (with nominally the same composition) in a helium environment (see

Figure 3.10). As much as a factor of ten increase in rupture life was observed for a given

applied stress for different heats of the alloy (Ennis et al. 1984).

Figure 3.11 shows the variation in creep rupture time with applied stress for Alloy 617 in

air and helium at several temperatures, compiled by Ennis et al. (1984). The helium used in

these experiments had a nominal composition (in Pa) of H2:CO:CH4:H2O in the ratio of

50:5:5:0.15, except for those at 1000°C, in which the environment was decarburizing to the

alloy. The few data points obtained at 1000°C under decarburizing conditions show a

significant decrease in rupture life when compared with the data developed in air. Figure 3.12

shows data from another study conducted in helium that was decarburizing to Alloy 617 at

950°C (Huchtemann 1989). The impurity levels in helium in this experiment was (in Pa)
43

Figure 3.11 Variation in creep rupture

time as a function of applied stress for

Alloy 617 in air and helium at several

temperatures (Ennis et al. 1984).

Figure 3.12 Variation in creep rupture

time as a function of applied stress for

Alloy 617 at 950°C in air and helium of

low steam content (Huchtemann 1989).

H2:CO:CH4:H2O:N2 in the ratio of 50:1.5:2:0.015:0.5. The data show a significant decrease in

creep life in helium when compared with that in air, for a given applied stress. Detailed

examination of the tested creep specimens showed that Alloy 617 did not develop a protective

oxide scale but exhibited significant internal oxidation of aluminum in the alloy accompanied

by carbon loss. Note that the helium environment in the gas-turbine-based systems will be

closer to the latter environment than those used in the study of either the steam-cycle system

or the PNP.

Figure 3.13 shows the stress dependence of time to accumulate 1% strain and of time to

rupture for Alloy 617 tested at 800 to 1000°C in air and in helium with “nominal” impurity

levels (Schubert et al. 1983). The data, developed up to 20,000 h, indicate that the chemistry

of the helium environment used had no significant effect on both the 1% strain and creep

rupture properties, although the data set exhibits significant scatter.
44

Figure 3.13 Time-to-1% strain and time-to-rupture as a function of applied stress for Alloy

617 in air and in HTGR helium (Schubert et al. 1983).

(a) (b)

Figure 3.14 Creep rupture ductility of Alloy 617 in air and helium: (a) Schubert et al. 1983

and (b) Cook 1984.

One of the interesting creep properties to compare is the rupture ductility for Alloy 617

that has been tested in air and in helium. Figure 3.14 shows a plot of elongation at rupture

from creep tests as a function of test time and temperature, obtained in two studies (Schubert

et al. 1983, Cook 1984). The data indicate that the alloy had much lower ductility in helium

compared to that obtained in air. Reason for these differences can be seen from the creep
45

Figure 3.15 Examples of creep curves and cracking behavior for Alloy 617 tested in air and in

helium at 950°C (Cook 1984).

curves and microstructures in Fig. 3.15, where air testing apparently induced surface cracks

early in the test, leading to higher initial creep rates (Cook 1984). In helium, the shape of the

creep curve indicated nucleation of surface cracks to be more difficult and virtual absence of

“tertiary” creep. Rupture was reported without development of a large number of creep cracks

in the helium environment. The apparent improvement of ductility in air was attributed to a

combination of oxidation and crack growth.

3.3.2 Fatigue Properties

Several low- and high-cycle fatigue studies have been conducted on Alloys 800H, 617,

Hastelloy X, and Hastelloy XR (Meurer et al. 1984, Soo and Sabatini 1984, Tsuji and Kondo

1984, Strizak et al. 1982, Kitagawa et al. 1979).

Strain-controlled fully reversed fatigue tests were conducted from room temperature to

871°C and at a cyclic strain rate of 4 x 10-3/s (Strizak et al. 1982). The helium environment

had a composition of (in Pa) H2:CO:CH4:H2O in a ratio of 30:2:3:0.2. Figures 3.16 and 3.17

compare data generated for Alloy 617 specimens in the solution-annealed and solutionannealed-plus-aged conditions, in which the aging and test temperatures were identical. The

results on thermal aging showed that aging for 10,000 h at these temperatures before testing

reduced fatigue life, but aging for 20,000 h improved fatigue life slightly. These changes were

attributed to microstructural changes that were noted in posttest examination of specimens.

Results of testing Alloy 617 in impure helium indicated an improved cycle life when compared

with the life obtained in pure helium. Strizak et al. (1982) also compared the data generated in

impure helium with those generated in air and concluded that the helium environment was in

no case detrimental to fatigue life for the alloy but was usually beneficial. Figure 3.18 shows a

comparison of the low-cycle fatigue behavior of Alloy 617 in air and helium with the design

curves (Schubert 1984).
46

Kitagawa et al. (1979) evaluated the low-cycle fatigue and creep fatigue behavior of

Alloy 617 at 1000°C in air and in helium. Figure 3.19 shows the low-cycle fatigue behavior of

the alloy in air and in 99.99% pure helium. Kitagawa et al. also tested some precarburized

specimens, the results for which are also shown in Figure 3.19. They concluded that the

fatigue properties in four-nine helium improved due to decarburization of the alloy, whereas

the precarburized specimens exhibited somewhat lower fatigue properties when compared with

data obtained in air. Figure 3.20 shows the fatigue lives of Alloy 617 under strain hold

conditions tested at 1000°C in helium. The fatigue lives decreased with an increase in strain

hold time.

Figure 3.21 compares experimental data on creep fatigue of Alloy 617 with the linear

summation rule for the creep and fatigue damage. It is evident that the damage under creepfatigue conditions is much more than predicted based on the linear rule. Additional data

under different loading scenarios are needed to develop a predictive capability on creep-fatigue

damage in the alloy, especially for helium purity levels typical of gas-turbine-based HTGRs.

Strizak et al. (1982) also developed low-cycle fatigue data on Hastelloy X in solutionannealed and solution-annealed-plus-aged conditions in air and in impure helium. The results

were similar to those observed for the Alloy 617 (see Figure 3.22).

Figure 3.16 Comparison of low-cycle

fatigue behavior of Inconel 617 tested in

air and in helium at 704°C (Strizak et al.

1982).
47

Figure 3.17 Comparison of low-cycle

fatigue behavior of Inconel 617 tested in

air and in helium at 871°C (Strizak et al.

1982).

Figure 3.18 Comparison of the low-cycle

fatigue behavior of Alloy 617 in air and

in helium with the design curves

(Schubert et al. 1983).

Figure 3.19 Comparison of low-cycle

fatigue behavior of solution-treated and

carburized Alloy 617 tested in air and in

helium at 1000°C (Kitagawa et al. 1979).
48

Figure 3.20 Influence of hold time on

fatigue life of Alloy 617 at 1000°C

(Kitagawa et al. 1979).

Figure 3.21 Comparison of experimental

creep-fatigue data with linear damage

accumulation rule (Schubert et al.

1983).
49

Figure 3.22 Comparison of straincontrolled fatigue data generated at

871°C in air and in helium for

Hastelloy X in several conditions

(Strizak et al. 1982).

3.3.3 Other Properties

Studies on creep and fatigue crack growth have been conducted with Nimonic 86

(Krompholz et al. 1984, Shahinian and Sadananda 1989).

Crack growth in nickel-base alloys IN-738 at 927°C under cyclic and static loads was

investigated using fracture-mechanics-type specimens (Shahinian and Sadananda 1989). The

influence of environment (specifically air, vacuum, and impure helium) on crack growth was

examined. Crack growth was markedly slower on a time basis under static load (creep) than

under cyclic load (fatigue), and in creep, crack growth required higher stress intensity for

initiation. Generally, the resistance to crack growth in creep was faster in air than in vacuum,

whereas in fatigue the converse was the case. The creep crack growth rates in helium

containing small amounts of H2, CO, and CH4 were not significantly different from rates

observed in vacuum.
50

4 Corrosion-Erosion

In systems that involve flow of particle-laden gases, one needs to consider corrosionerosion as a mode of material degradation. Such information is needed to assess the long-term

performance of candidate materials. Erosion, by itself, is a complex phenomenon and can be

influenced by several variables, and combined erosion-corrosion is even more complex because

of the dynamic changes that occur on the eroding surface due to gas-metal interactions.

The main operational experience with a gas stream that is laden with carbon-dust

particles comes from the 15-MWe, Arbeitsgemeinshaft Versuchsreaktor (AVR), a pebble-bedtype HTGR. The reactor operated for 20 y and contained 100,000 fuel balls. At the end of

20-y operation, 60 kg of carbon dust was collected. This corresponds to 0.03 g/ball·y.

Assuming the carbon dust particles are 1 μm dia, the number of particles created would be

2 x 109 particles/ball·y. If one assumes an average helium inventory in the reactor of 10 m3,

the particle density in the gas stream would be 2 x 1013 particles/m3·y, or 4 x 1014

particles/m3 at the end of 20-year operation or end of reactor life. In a closed-loop cycle, the

helium replenishment is of the order of 0.1%/day, which means the entire helium inventory

changes approximately every 3 years, and the dust loading increases with the aging of the

reactor. The AVR experience shows that the blowers were significantly degraded over the

years, and the cause was not reported. One of the causes for the dust in the AVR was

attributed to an excessive release during control rod insertions into the reactor core. In the

proposed current design with separate channels for coolant and control rod insertions, the

dust loading may be less than that in AVR.

The PBMR-type reactor will contain on the order of 1.2 million fuel balls, and the helium

inventory will be correspondingly high. However, the particle density will probably be less than

that for the AVR since special channels, with no helium flow, are proposed for the control rod

insertions. From the erosion standpoint, the consequences of particle flux, temperature, and

particle size on the structural materials and on direct-driven gas-turbine materials are not

known.

From the erosion-corrosion standpoint, degradation of materials is strongly influenced by

factors such as erodent hardness, erodent size, temperature, substrate mechanical properties,

and particle flux and velocity (Majumdar et al. 1988). An assessment of the erosion

performance of structural alloys with different erodents showed that the erosion rate increases

with the hardness of the erodent. Carbon dust particles with an average Vickers hardness of

20 and an average particle size of 1 μm (assuming no agglomeration) will tend to deposit on the

surfaces of components rather than cause erosive wear.

The deposits on the surfaces of structural components and power conversion components

alter the chemistry and thereby disrupt the protective capacity of the scales, if any, against

corrosion. Liu and Natesan (1988) have examined methodologies for assessing and predicting

metal oxidation-vaporization-erosion. The study examined the synergistic interaction of

oxidation and erosion of metals and alloys by considering three competing processes: oxide

scale thickening via growth, thinning via vaporization, and erosion. The authors demonstrated

that not only do the models numerically reproduce the general features of oxidation-erosion

observed in the few reported experiments on pure metals and alloys, but they also provide

considerable insight into the conditions under which one mode (oxidation or erosion) can
51

become dominant over the other. Such a corrosion-erosion analysis of the reactor system and

the power conversion system may be useful in evaluating and quantifying the issue and may

also shed light on the purification scheme needed for the helium supply and replenishment

and chemistry control.
52

5 Design Methodology

A companion report (Shah et al. 2003) reviews and evaluates currently available national

and international codes and procedures for use in the design of HTGRs, including, but not

limited to, the GT-MHR design and the PMBR. That report evaluates the applicability of the

codes, standards, and procedures for the metallic materials that have been used or

recommended for HTGRs, taking into account the HTGR operating temperature and

environments. Seven codes and procedures [five ASME Codes and Code Cases, one French

Code (RCC-MR), and one British Procedure] were reviewed and evaluated. The ASME Codes

and Code Cases included Section III, Subsection NB and Subsection NH; Code Cases N-499-1,

N-201-4; and the draft Code Case for Alloy 617. Two major findings were obtained. (1) The

maximum temperature permitted by the codes and code cases for the materials acceptable for

HTGR components, except for those made of Alloy 671, is lower (760°C) than the maximum

temperature (850°C or higher) that these components may experience in reactor service. The

scope of the code and code cases needs to be expanded to include materials with allowable

temperatures of 850°C. (2) The codes and code cases do not provide specific guidelines for

environmental effects, especially the effect of impure helium, on the high temperature behavior

(e.g., creep and creep-fatigue) of the materials considered. The report concluded that needed

data on environmental effects should be collected or generated, if not available, so that specific

guidelines can be developed for these effects.
53

6 Off-normal Situations

Several off-normal situations can arise and lead to changes in the environment within the

reactor and thereby can impact the metallic components, such as reactor internals, piping,

power conversion systems, etc. These include

• Loss of coolant

• Air and/or steam ingress

• Particulate-laden gases

6.1 Loss of Coolant

The loss of coolant can increase the temperature of the reactor core, thereby causing

failure of fuel pellets and release of fission products. Fission products such as Cs, Sr, and I

can affect the structural materials, depending on how much is released and for how long the

release was present during this off-normal event. However, a system such as GT-MHR has

three main protective barriers to confine fission products (IAEA 2001). Figure 6.1 show a

schematic for fission-product barriers and retention properties of TRISO fuel elements for a

typical HTGR. The barriers include:

• The fuel kernel: The fuel is effective in confining solid fission products and partially

retains gaseous and volatile fission products. The fuel temperature level and burnup

are limited in design to enable confinement of fission products within the fuel kernel.

• Fuel kernel coatings: The next and most important barrier for fission product

confinement is the protective coating on each fuel kernel. The GT-MHR design uses

fuel with coatings made from three layers of pyrolitic carbon of different density and

one layer of SiC. Experimental studies with coated particles of this type have

validated their capability to confine fission products up to 1600°C, which corresponds

to the maximum possible temperature level under emergency conditions.

• Graphite: The graphite of the fuel compacts and fuel blocks where the fuel particles

are located can act as a barrier to confine fission products (such as Cs, Rb, Se, Ba,

rare-earth elements) fairly effectively.

In addition, the fraction of particle failures during manufacture had a specification of

6 x 10-5; that is, on average, one particle with a defective SiC layer is to be expected in two to

three fuel spheres (Kroger et al. 1988). Also, TRISO coatings do not start to thermally degrade

until temperatures approach 2000°C (LaBar 2002). Normal operating temperatures do not

exceed about 1250°C. Even in the worst case, temperatures can be maintained below 1600°C

(see Figure 6.2). Based on this analysis, a minimal release of fission products is expected, and

its effect on the metallic structural materials is probably negligible.
54

Figure 6.1 Fission-product barriers and retention properties of TRISO fuel elements for a

pebble-bed gas reactor: (a) particle and (b) fuel element (Kroger et al. 1988).

Figure 6.2 Coated-particle fuel temperature capability (La Bar 2002).

6.2 Air and Water Ingress

The potential ingress of air or steam can lead to a significant reaction with the graphite in

the core and possibly with the fuel coatings. Such a reaction would significantly change the
55

chemistry of the helium coolant. In turn, this change can impact the reliability and integrity of

metallic components. Katscher and Moormann (1986) have analyzed the kinetics of reactions

between the structural graphite and the oxygen, carbon dioxide, and water vapor that are

present due to accidental ingress. They calculated maximum burnoff rates for graphite with an

unlimited supply of oxygen and concluded that there is uncertainty in the calculations

influenced by the size and location of the leaks, ingress rate of air, mixing processes, etc.

The GT-MHR design includes technical measures that make the ingress of air or steam

into the core less likely and also allow for decreasing the consequences should this event occur

(IAEA 2001). These measures include:

• The working pressure of water in the PCS coolers and helium/water heat exchanger is

considerably less than the helium working pressure in the primary circuit ( 7 MPa in this

circuit versus 0.35 to 0.80 MPa in the intermediate water circuit in PCS coolers). This

nearly eliminates the possibility of water ingress into the core during reactor power

operation or cool down, if pressurized conditions are maintained in the primary circuit.

• Implementation of the “leak before break” concept limits the scope of a potential loss-ofintegrity event in the vessel system. The leak-tight containment minimizes accidental

ingress of air into the reactor. In the event of air leakage, helium is replaced by air over a

long duration due to mass exchange and diffusion processes. The amount of air flow would

be small through the place of leakage. Gases generated by the graphite oxidation process

prevent the air coming to the core graphite structures, thereby moderating the oxidation

process. The graphite temperature decreases during the reactor cool down to 600°C,

which leads to shutting down of the oxidation process. It was estimated that not more than

1% of the total graphite mass would be subjected to corrosion. The impact of helium with

such impurity levels, albeit for short times, on the metallic components is not established.

6.3 Particulate-Laden Gases

One of the purposes of the helium services system (HSS) in the GT-MHR is to maintain

the primary coolant composition at specified levels to provide necessary conditions for longterm and reliable operation of the primary circuit components at normal operation and in

accidents associated with water or air ingress into the primary circuit (IAEA 2001). The HSS

includes the helium purification system (HPS), which consists of the equipment for helium

cleanup of chemical and radioactive impurities, heat exchange equipment, pipelines, valves,

instrumentation sensors, and the means for regeneration.

The HPS ensures the following levels of impurities in the helium under normal operating

conditions: H2O < 2 vppm, CO2 + CO < 6 vppm, H2 < 5 vppm. The primary helium flow rate

through the HPS is 0.5 kg/s. The total helium flow rate through the core is 316 kg/s. This

rate indicates that only about 0.15% of the reactor helium is replenished at any given time.

Therefore, particulates such as carbon dust can accumulate in the helium stream as a function

of time, even though the amount is not established. Furthermore, the role of the carbon dust,

if any, on the deposition and/or erosion of metallic structural materials and PCS materials is

not known.

Even though the release of fission products during normal operation is projected to be

negligible, Kroger et al. (1988) reported that based on AVR studies, 73% of the released cesium
56

would be deposited on cooler metallic surfaces, 5% would be borne by graphite dust, and the

rest would be retained by the reflector graphite. The role of such a deposition on the metallic

surfaces in the degradation, if any, of metallic components is not established. It has been

mentioned that release of the fission products deposited in the primary circuit outside the core

is possible in certain accident situations. In the case of reactor depressurization, dust-borne

activity deposited in corners of the streaming channel or on surfaces exposed to flow can be

remobilized by an increased flow velocity. The impact of these scenarios on the

materials/components integrity is not known.
57

7 Summary and Recommendations

The advanced HTGR concept, currently considered by the Member States of the IAEA’s

International Working Group on Gas Cooled Reactors, involves helium as the coolant and a

closed-cycle gas turbine for power generation with a core outlet/gas turbine inlet temperature

of 850°C and a net plant efficiency of 47%. In these designs, the helium from the core is used

directly to drive the turbines or indirectly to heat the air or nitrogen, which drives the turbines.

The system concepts include the GT-MHR developed by a consortium led by General Atomics in

the U.S., the PBMR designed by ESKOM of South Africa and British Nuclear Fuels of U.K., and

the HTTR of Japan.

This report has reviewed and evaluated the available information on performance and on

long-term behavior of materials in environments that are typical of helium-cooled gas reactors.

As a part of this evaluation, we have identified the materials that have been used in HTGRs,

their functions, and the environmental conditions under which they have operated.

We likewise identified the materials to be used in new HTGRs, the functions they will serve,

and the anticipated environmental conditions under which they are expected to operate. We

evaluated those HTGR conditions that may affect material property response, such as the

steady-state operating environment, component geometry, and off-normal situations. Several

conclusions were drawn from this evaluation:

• The majority of materials research and development programs in support of HTGRs

was conducted in the 1960s to early 1980s. The thrust of these programs was to

develop a database on materials for application in steam-cycle and process-nuclearheat based HTGRs.

• Very little work was done on materials with emphasis on direct and/or indirect

gas-turbine-based HTGRs during this period.

• The primary materials for high temperature application in HTGRs that have been

evaluated in detail are Alloys 800H and 617 and Hastelloy X. Among them, Alloy

800H is code certified for temperatures up to 760°C for use in nuclear systems.

A draft code case for Alloy 617 has been developed. A substantial database has been

developed for both Alloys 800H and 617, and a limited database exists for

Hastelloy X. Since the high temperature scaling in Hastelloy X has not been

adequate, a modified version Hastelloy XR has been developed in programs conducted

in Japan.

• Extensive analysis was made on the influence of helium coolant on the chemical

compatibility of structural materials that are planned for use in HTGRs. Even though

helium, by itself, is inert towards the materials, it is often contaminated by small

amounts of gaseous impurities such as H2, H2O, CH4, CO, CO2, and O2 from a

variety of sources in the reactor circuit. The gas chemistry and the thermodynamic

activity for carbon and oxygen in the gas phase are difficult to ascertain because of

the nonequilibrium nature of the gas mixture. Furthermore, most studies on gas

chemistry simulations were performed close to atmospheric pressure, whereas the

system pressure in the reactor is 7 MPa.
58

• Structural alloys can be significantly corroded by the gaseous impurities in helium at

elevated temperatures. Past studies have shown that the corrosion of heat-resistant

materials such as austenitic stainless steels, Alloy 800H, and Alloy 617 may involve

oxidation, carburization, and decarburization, depending on the exposure

temperature, carbon activity in the gas phase, and the alloy composition.

Furthermore, the corrosion process is “dynamic” in that it is dictated by the exposure

time, gas chemistry variations, integrity of the corrosion product scales, and presence

of particulates in the gas phase. A thermodynamic framework that includes

evaluation of chemistry under nonequilibrium conditions is presented in this report.

The available corrosion information for several candidate alloys exposed to helium

environments is evaluated to address the scaling behavior and its relationship to gas

chemistry.

• Available data are reviewed on the mechanical properties of candidate alloys in a

helium environment with low levels of impurities. In most of these tests the

environment contained sufficiently high H2O (typical of steam-cycle-based HTGRs),

and the alloys predominantly developed oxide scales on the surface during

mechanical testing. The mechanical properties evaluated included creep, creep

rupture, low- and high-cycle fatigue, creep fatigue, and fracture toughness. The

analysis showed that the damage under creep-fatigue conditions is much more than

prediction based on the linear rule. Additional data on creep-fatigue under different

loading scenarios are needed to develop a predictive capability on creep fatigue

damage in the alloy, especially in helium purity levels typical of gas-turbine-based

HTGRs.

• In studies that were conducted in a helium environment with reduced levels of H2O,

Alloy 617 did not develop protective oxide scales, and the creep properties of the alloy

showed a decrease in life when compared with data generated in air. The aluminum

content in Alloy 617 led to internal oxidation rather than an external scale, especially

in environments with low oxygen partial pressure. It is believed that the impurity

levels in gas-turbine based systems will be predominantly decarburizing since oxygen

potential is expected to be much lower than that in steam-cycle based systems.

• Before materials performance can be assessed, it is essential to establish an

operational window in terms of impurity levels and particulate concentrations in

helium that can be maintained and controlled with certainty. It was demonstrated

that not only do the corrosion and erosion models numerically reproduce the general

features of oxidation and/or erosion observed in the few reported experiments on pure

metals and alloys, but they also provide considerable insight into the conditions

under which one mode (oxidation or erosion) can become dominant over the other.

Such a corrosion-erosion analysis of the reactor system and the power conversion

system may be useful in evaluating and quantifying the issue and may also shed light

on the purification scheme needed for the helium supply and replenishment and

chemistry control.

• Over the last 20 years, alloy manufacturers have developed several alloys with

improved resistance to oxidation and to carburization/decarburization at elevated

temperatures. These materials include Alloy 230, 602CA, microalloyed cast HP alloys,

and W-containing Ni-base superalloys. It is beneficial to examine the scaling and
59

corrosion performance of these alloys (in addition to Alloys 617, Hastelloy X, etc.) in

helium with low oxygen partial pressures for their viability in the HTGRs.

• An extensive database is needed on candidate alloys for application in turbines (both

blades and disks), especially for direct-cycle systems.

• The off-normal situations addressed in this report are limited to those that can lead to

an increase in temperature and/or change in the chemistry of the environment, both

of which can impact the metallic components used within the reactor.

• The current status of the materials database for the pressure vessels, piping and

internals, and power conversion applications are summarized in Tables 7.1-7.3.
Table 7.1 Summary of available thermomechanical data for alloys considered for pressure vessel

and other non-core, lower temperature components

Alloy

Property

2-1/4Cr-1Mo P9 P91 P92 E 911 HT9 316 SS

Tensile Strength

Creep

Fatigue

Creep Fatigue

Long-Term Aging:

Effects on

Microstructure

Environmental

Effects

X X X X X X X

ASME Code

Approved

Yes Yes Yes Yes ? Yes Yes

√ = Data sufficient; X = Additional data needed, especially in He environment.

60
Table 7.2 Summary of available thermomechanical data for various alloys considered for piping and internals

Alloy Name

Property

800H Hastelloy X Inconel 617 Hastelloy XR Nimonic-86 Manaurite Inconel 519

Tensile Strength ASME code approved

for T < 760°C. Tensile

data available

T < 900°C

Limited

data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

1200°C

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Creep Limited data

available for

T < 760°C

Limited

data

available for

T < 1100°C

Limited data

available for

T < 1100°C

Limited data

available for

T < 1100°C

Limited data

available for

T < 1100°C

Limited data

available for

T < 1100°C

Limited data

available for

T < 1100°C

Fatigue Limited data

available for

T < 760°C

Limited

data

available for

T < 1100°C

Limited data

available for

T < 1100°C

D Limited data

available for

T < 1100°C

Limited data

available for

for

T < 1100°C

Limited data

available for

for

T < 1100°C

Limited data

available for

for

T < 1100°C

Creep Fatigue Limited data

available for

T < 760°C

Limited

data

available for

T < 760°C

Limited data

available for

T < 760°C

Limited data

available for

T < 760°C

Limited data

available for

T < 760°C

Limited data

available for

T < 760°C

Limited data

available for

T < 760°C

Aging Limited data

available for

T < 1200°C

Limited

data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Limited data

available for

T < 1200°C

Environmental:

Effects, especially in

He with controlled

impurity

Susceptible to metal

dusting at

400°C < T < 800°C

Need to

generate

data for

creep,

fatigue and

creepfatigue

Need to

generate data

Need to generate

data

Need to

generate

data

Need to

generate

data

Need to

generate

data

Code Case Status Approved N-201-4

T < 760°C

Not

approved

Draft ASME

code case at

T < 982°C

Not approved Not approved Not approved Not approved

61
Table 7.3 Summary of available thermomechanical data for various alloys for power conversion applications*

Alloy

Property

713LC MAR-M 004 M21 738 Udimet 520 Mo-TZM A286 706 718 Waspaloy Udimet 720 MA6000

Tensile T 900°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C T 1200°C

Creep T °C T 1100°C T 1100°C

Fatigue T °C T 1100°C T 1100°C

Creep Fatigue X X X

Aging X X X

Environmental

effects

X X X X X X X X X X X X

Code case

status

NA NA NA NA NA NA NA NA NA NA NA NA

 = Limited data available; X = Need to generate data and develop phenomenological performance-based models; NA = Not approved.

*Power conversion requires high temperatures. Currently, no ASME code case covers the temperature range of interest ( 900-1100°C).

62
63

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