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Crack resistance of age-hardening Fe-Ni alloys in gaseous hydrogen

Conference Paper · January 2010

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 1

CRACK RESISTANCE OF AGE-HARDENING Fe-Ni ALLOYS

IN GASEOUS HYDROGEN

A.I. Balitskii 1), L.M. Ivaskevich1), V.M. Mochulskyi1)

1)
Karpenko Physico-Mechanical Institute of National Academy of Sciences
of Ukraine, Naukova str.,5, UA 79601, Lviv, Ukraine

ABSTRACT

The effect hydrogen on short-term strength, low-cycle durability, fatigue threshold ΔKth,
fatigue fracture toughness Kfc and fracture toughness Kc of Fe-27Ni and Fe-55Ni alloys in
wide diapason of pressure and temperature was investigated. The parameters of loading
and the modes of hydrogen action for which the crack resistance parameters of alloys are
minimum was established. Fatigue threshold ΔKth, fatigue fracture toughness Kfc and
fracture toughness Kc of Fe-27Ni steels is decreased by hydrogen in a wide range of
pressures and temperatures. The plane-strain state conditions and, correspondingly, the
conditions of the evaluation of KІс were fulfilled after preliminary high-temperature
hydrogenation of compact tension specimens with a thickness of 20 mm at hydrogen
concentration 8 ppm in the temperature range 293…483 K. Fracture toughness Kc of Fe-
55Ni alloys was decreased at 293 K from 116 MPa√m in helium to 78 MPa√m in hydrogen
under the pressure 30 MPa and to 68 MPa√m in hydrogen under the pressure 30 MPa after
preliminary high-temperature hydrogenation (hydrogen concentration 15 ppm). The plane-
strain state conditions were not fulfilled at both case. Low-cycle durability are most sensitive
to the hydrogen degradation (decrease by 80 and 70% of the value in helium, respectively).
The parameters of cyclic loading ΔKth and Kfc decreased by 25%, fracture toughness Kс
and reduction in area ψ under short-term static tension – by 50%

KEYWORDS

Hydrogen resistance, fracture toughness, short-term strength, low-cycle durability,

fatigue threshold, fractographic peculiarities

INTRODUCTION

The production of turbine blades and other parts of air-plane and rocket engines, nuclear
reactors, petrochemical equipment requires the wide usage of dispersively hardened heat-
resistant alloys. Particularly this concerns Fe-Ni alloys. In these products heat-resistant alloys
are exploited at high temperatures in the contact with high-pressure hydrogen. Therefore one
of the most important requirements for such alloys is their resistance to hydrogen degradation,
in other words their ability to keep high level of mechanical properties under the action of
hydrogen in wide range of exploitation parameters. At the same time, age-hardening alloys
are known to be rather sensitive to hydrogen embrittlement [1, 2, 3].

The aim of this investigation is to study the influence of high-pressure hydrogen on the short-
term strength, low-cycle durability, static and cyclic crack resistance of
10Kh15Ni27Ti3W2MoAl and 05Kh19Ni55Nb2Mo9Al alloys in the temperature range 293-
1073°K.
 2

MATERIALS AND TEST PROCEDURE

The chemical composition, heart-treatment modes and original properties are given in Tables
1 and 2.

Table 1. Chemical composition of Steels

Content of elements, wt.%

Alloy
С Si Cr Ni Мо W Ti Al Fe
10Kh15Ni27Ti3W2MoAl 0.09 0.6 15.18 27.11 1.41 1.92 2.85 0.29 Bal
05Kh19Ni55Nb2Mo9Al 0.05 0.23 19.04 Bal 8.87 - - 1.49 12.0

Table 2. Modes of Thermal Treatment and Mechanical Properties of Steels in

Helium/Hydrogen (35 MPa) at Room Temperature

Thermal Treatment Mechanical Properties

Number cyc-
Alloy Solution Mode of σu, σ0,2 δ ψ les to fractu-
treatment aging MPa MPa % % re, bending
strain 1.6 %
1023К, 16 h 1270 870 17 23 2277
10Kh15Ni27Ti3W2MoAl 1373К, 1 h
923К, 10 h 1240 880 15 20 1667
1000К, 15 h 1080 65 35 38 2560
05Kh19Ni55Nb2Mo9Al 1323К, 1 h
923К, 10 h 970 66 7 21 299

Static tensile tests were carried out on standard five-fold cylindrical specimens by
displacement rate V = 0.1 mm/mm. During the test the specimen was positioned into the
chamber specially designed for high-temperature tests at 293...1093 K temperature range
under 0.1...35 MPa hydrogen pressure. Specimens were tested in 293...1073 K temperature
range under 35 MPa hydrogen pressure and, as a comparison, in helium. The low-cycle
durability for pure strain-controlled sign-preserving bending was found under pressures of
35MPa for the amplitudes ε = 1.6% and a loading frequency of 0.5 Hz on polished plane
specimens with a working part of 3 x 6 x 20 mm. The stress intensity factor under static
loading Kc is computed either for the maximum force Fc in the "F-V" linear diagram or for the
force FQ determined by using the 5% secant for nonlinear diagrams [4]. Rectangular
compact specimens 50 x 60 x 20 mm in size were tested for eccentric tension in a high-
pressure chamber mounted on a UME-10TM tensile-testing machine under pressures of 0.4-
30 MPa at a strain rate of 0.1 mm/min. The values of Kc can be found by using the Srawley-
Gross formula [5]. The cyclic crack resistance (CCR) characteristics were determined in
three-point bending of 160x40x 20 mm beam specimens at a frequency of loading of 20 Hz
and a coefficient of cycle asymmetry R = 0.22 [6, 7]. A fatigue crack 2 mm long was
preliminarily grown from a stress concentrator. We placed specimens with a crack in a high-
pressure chamber, set on EUS-40 unit, where, upon attaining required parameters of the
hydrogen atmosphere, loaded them according to a known technique [7], and determined the
threshold ΔKth in hydrogen. The lengths of cracks were measured by changing the
parameters of the electromagnetic eddy field. Transducers were preliminarily calibrated in air
by comparing changes in eddy currents caused by the growth of a crack, and changes in the
length of cracks measured optically on a specimen with a KM-8 cathetometer. We
constructed kinetic fatigue fracture diagrams (KFFD) in helium and hydrogen under
pressures up to 10 MPa at temperatures of 293 to 673 K from experimentally obtained
lengths of cracks Δl during ΔN cycles under a load F*.
 3

To determine the indicated mechanical characteristics in hydrogen, the working chambers

were preliminarily evacuated, blown-out with hydrogen, again evacuated, and filled with
hydrogen up to a given pressure. At high temperatures, the specimens were held under the
testing conditions for 30 min up to the attainment of thermal equilibrium. It has been
established [1-3, 8, 9] that, at some values of hydrogen pressure and strain rate, which
depend on the chemical compositions and structures of materials, maximum influences of
hydrogen on the plasticity, low-cycle fatigue life, and static and cyclic crack resistance of
martensitic steels and nickel alloys are achieved. In short-term tension, austenitic dispersion-
hardened steels are substantially embrittled by hydrogen after preliminary hydrogenation at
elevated temperatures and upon attaining its content above 12 ppm, and the properties of
hydrogenation specimens at room temperature in air and hydrogen are equal [1, 8, 9]. This is
why we held a part of the specimens for 10 h in a hydrogen atmosphere under 623 K and 35
MPa. These regimes provide the hydrogenation of specimens to hydrogen contents of 15
ppm. Hydrogenated and non-hydrogenated specimens were tested in helium and hydrogen
under different pressures.

The hydrogen content in iron-nickel alloys was determined with a LECO TCH 600
instrument.

EXPERIMENTAL RESULTS

Influence of hydrogen pressure on the mechanical properties of alloys.

The dependence of number of cycles to fracture N (curve 5, 6 in Fig. 1b), relative elongation
δ (curve 1, 2 in Fig. 1b), reduction of area ψ (curve 3, 4 in Fig. 1b) and fracture toughness Kc
(curve 2, 4 in Fig. 2) of 05Kh19Ni55Nb2Mo9Al alloy on the hydrogen pressure consists of
two regions. In the first region (low pressure) the mechanical properties sharply drops when
pressure increased, and in the second, the negative action of hydrogen becomes stable.
This means that there exits a pressure under which hydrogen degradation of this material
reaches its limit. Additional influence of preliminarily dissolved hydrogen on the mechanical
properties of 05Kh19Ni55Nb2Mo9Al alloy is detected only for low pressures (Fig. 1b, Fig. 2).
N,cycles 120
Kc,MPa m

2000 N,cycles
2000 5
b
5 2
a
1000 6 100
1000 1
6

,% ,% 80 4
3 3
25 30
20 4
1
20 1 60
15 3
42 2
10 10
0 5 10 15 20 25 30 0 5 10 15 20 25 30 0 5 10 15 20 25 30
Р, M
Pa Р, MPa Р,MPа

Fig. 1. The relative elongation δ (1,2), reduction of area ψ Fig. 2. Fracture toughness Kc
(3,4) (V = 0,1 mm/min.) and number cycles to failure N of 10Kh15Ni27Ti3W2MoAl (1,
(5,6) (ε = 1,6%) specimens of 10Kh15Ni27Ti3W2MoAl (a) 3) and 05Kh19Ni55Nb2Mo9Al
and 05Kh19Ni55Nb2Mo9Al (b) alloys versus hydrogen (2, 4) alloys versus hydrogen
pressure P at 293 K: 1,3,5 – non-hydrogenated specimens; pressure P at 293 K: 1,2 –
2,4,6 – hydrogenated specimens. non-hydrogenated specimens;
3,4 – hydrogenated
specimens.

On the contrary at room temperature, hydrogen reduces N, δ, ψ and the parameter Kc of the
10Kh15Ni27Ti3W2MoAl steel in the entire range of pressures (Fig.1 a, curve 1, 3, 5; Fig. 2,
 4

curve 1), i.e., a pressure of 30 MPa is insufficient for a maximum decrease in mechanical
properties. However, the minimum values of this properties which are independent of
hydrogen pressure, are attained after preliminary high-temperature hydrogenation up to
concentration 15 ppm (Fig.1 a, curve 2, 4, 6; Fig. 2, curve 3). Hydrogen decreases the
coefficient of crack resistance and effects on the character of fracture. Under the conditions
of maximum hydrogen embrittlement, the load displacement diagrams become linear with
sharp maxima (as functions of load) and correspondent to type II [4]. The values of Kc can be
regarded as equal to K1c, i.e., they satisfy the condition l, b ≥ 2.5(KcH/σ0,2H)2, where l is the
crack length and b is the thickness of specimen [4]. The indicated condition is satisfied at
room temperature for 10Kh15N27T3B2MR steels after preliminary hydrogenation (623 K, 35
MPa, 10 h) for all hydrogen pressure with K1c is equal 50–55 MPa√m. Fracture toughness Kc
of 05Kh19Ni55Nb2Mo9Al alloys was decreased at 293 K from 116 MPa√m in helium to 78
MPa√m in hydrogen under the pressure 30 MPa and to 68 MPa√m in hydrogen under the
pressure 30 MPa after preliminary high-temperature hydrogenation (hydrogen concentration
15 ppm). The plane-strain state conditions were not fulfilled at both case.

Influence of temperature on the hydrogen degradation of alloys under short-term

static and low-cycle loads.

As temperature increases from 293 to 1073 K, the ultimate strength σu both alloys in
hydrogen and helium decreases (Fig. 3). The strength of 05Kh19Ni55Nb2Mo9Al alloy is
much lower than the strength of 10Kh15N27T3B2MR steel in whole temperature range (Fig.
3, curve 1-4). Plasticity characteristics δ and ψ are much more sensitive to the influence of
hydrogen then ultimate strength σu (Fig. 3, curve 5-10). The strength of alloys significantly
decreases in neutral media at temperatures higher than 873 K. This process is accompanied by
a decrease in the plasticity of both materials at 973°K and a significant increase in the relative
elongation and reduction of area for specimens made of 05Kh19Ni55Nb2Mo9Al alloy at 1073 K
(Fig.3, curve 5-7). The drops of plasticity of the dispersion-hardening austenitic alloys in the
temperature range of intense phase transformations are caused by the localization of strains on
the grain boundaries due to the intense redistribution of Ni, Ti, and Al in the boundary regions.
Moreover, the increase in plasticity observed at higher temperatures is caused both by the partial
coagulation of hardening phases and possible dissolution of small amounts of finely divided
precipitations [10].

1200
u, MPa Fig. 3. Temperature dependences of
1000
3 4 ultimate strength σu (1-4) and
800 1 2
reduction of area ψ (5-8) of
600 10Kh15Ni27Ti3W2MoAl (1, 2, 8-10)
and 05Kh19Ni55Nb2Mo9Al (3, 4, 5-7)
80
 alloys (V = 0,1 mm/min.) at helium (1,
60 3, 5, 8) and hydrogen (2, 4, 6, 9) under
5
40 the pressure 30 MPa: 1-4, 5, 6, 8, 9 –
6 7
20 8 non-hydrogenated specimens; 7, 10 –
0 9 10 hydrogenated specimens.
400 600 800 T, K 1000

As temperature increases, the degree of hydrogen embrittlement of the non-hydrogenated

specimens made 10Kh15Ni27Ti3W2MoAl steel varies insignificantly (Fig. 3, curve 9),
whereas the degree of embrittlement of preliminarily hydrogenated specimens decreases
(Fig. 3, curve 10). As a result, at temperatures higher than 973°K (when the amount of
hydrogen sufficient for the maximum hydrogen degradation under the conditions of short-
term static tension penetrates into the metal in the course of the experiment), the levels of
 5

plasticity of the hydrogenated and non-hydrogenated specimens in hydrogen-containing

media are identical (Fig. 3, curves 9 and 10). The action of hydrogen on the plasticity of
05Kh19Ni55Nb2Mo9Al alloy is maximum at room temperatures, minimum at 873 K, and again
noticeable at 1073 K (Fig. 3, curves 5-7). As temperature increases from 293 to 473 K, the
influence of hydrogen on the low-cycle durability of non-hydrogenated specimens of
10Kh15Ni27Ti3W2MoAl steel decreases from 0.86 to 0.22 (Fig. 4a, curve 4). At the same,
for hydrogenated specimens, it remains constant within the range 0.2-0.22 up to 773°K (Fig.
4a, curve 5). At 473 K, the additional influence of preliminarily dissolved hydrogen becomes
negligible (Fig. 4a, curves 4 and 5), i.e., for low-cycle fatigue, this temperature is sufficient for
the hydrogenation of dispersion-hardening austenitic steels from the atmosphere of
hydrogen. In tension, when the maximum tensile stresses are localized at much larger
distances from the specimen surface, this effect is attained for 973 K (Fig. 3, curves 9, 10).
The degree of embrittlement of 10Kh15Ni27Ti3W2MoAl steel somewhat decreases at 873 K.
However, for both loading modes, the temperature interval of significant hydrogen
degradation for austenitic steels is much larger than for 05Kh19Ni55Nb2Mo9Al alloy (Figs. 3
and 4a, b).

N, cycles
N,cycles

2000 1
1 a b
2000 2
2 3
1000 1000
3
N,
N,

0.9 0.9 4
5
0.6 4 0.6

0.3 5 0.3
0.0 0.0
400 600 800 T, K 1000 400 600 800 T, K
Fig. 4. Temperature dependences of number cycles to failure N (1-3) (ε = 1,6%) and
coefficients of hydrogen influence βN = NH/NHe (4, 5) on 10Kh15Ni27Ti3W2MoAl (a) and
05Kh19Ni55Nb2Mo9Al (b): 1, 2, 4 – non-hydrogenated specimens; 3, 5 – hydrogenated
specimens.

Influence of temperature on the static and cyclic crack resistance in hydrogen.

The character of influence of the temperature of tests depends on their conditions (Fig. 5).
With increase in temperature, the crack resistance in helium first rises from 98 (293 K) to 120
MPa√m (423 K) and, in the range 423-693 K, hardly changes (curve 1); in this case, the
thickness of the specimens is insufficient for the realization of the plane strained state.
Gaseous hydrogen reduces noticeably the crack resistance even at room temperature, and
this effect is dramatically enhanced as the temperature increases to 423 K (curve 2). In the
range of rather low temperature, preliminary hydrogenation influences most, and, at 423 K,
the crack resistances in hydrogen of preliminarily hydrogenated (curve 3) and non-
hydrogenated (curve 2) specimens are equal, i.e., at this temperature, a sufficient amount of
hydrogen penetrates in the material from the gas phase even during the experiment. As the
temperature further rises, the negative influence of hydrogen is weakened in the absence of
difference in the values of Kc between the hydrogenated and non-hydrogenated specimens,
but, even at a maximum temperature, the parameter Kc in hydrogen is 21% smaller than that
in helium.

The hydrogenation of the material changes the character of loading-displacement diagrams,

which describe the plane stressed state in crack resistance tests [11]. Under a maximum
hydrogen embrittlement (after preliminary hydrogenation at all hydrogen pressures), they
become linear, and fractures of specimens are covered with cleavage facets, which indicates
 6

that the conditions of plane deformation are satisfied and enables us to determine the value
of Klc. With decrease in the rate of tension of the hydrogenated specimens from 0.5 to 0.1
mm/min, the crack resistance does not decrease additionally. This is why we can state that
the values of KIc of 10Kh15Ni27Ti3W2MoAl steel are minimum under the action of hydrogen.

120 60
Kc,MPam

Kfc, MPam
110 3
50
100 1

90
4
40
80 2
10 1

Кth,
70
8
60
2
3 6
50 4
300 400 500 600 700 300 400 500 600 Т, К 700
T,K
Fig. 5. Temperature dependence of static Fig. 6. Temperature dependence of
stress intensity factor Кс fatigue threshold ΔKth (1, 2) and fatigue
10Kh15Ni27Ti3W2MoAl steel in helium (1), in fracture toughness Kfc (3, 4) in helium (1,
hydrogen under pressure 10 МPа (2) and in 3) and in hydrogen under the pressure 10
hydrogen under pressure 10 МPа after hydro- МPа after hydrogenation (623 K;10
genation (673 К;10 МPа;10 h.) (3). МPа;10 h) (2, 4).

The temperature range for which the sizes of specimens are sufficient for the determination
of Klc under maximum hydrogen degradation is 293-483 K. As the testing temperature rises,
the values of ΔKth of the alloys in helium and hydrogen increase, whereas the characteristic
Kfc is practically independent of temperature in both atmospheres (Fig. 6). The influence of
hydrogen is maximum at 293 K (ΔKth diminishes from 6.81 MPa√m in tests in helium to 4.95
MPa√m for hydrogenated specimens and tested in hydrogen under a pressure of 10 MPa
(curves 1, 2). The decrease of hydrogen concentration to 8 ppm does not diminish the
negative effect, i.e., the minimum values of ΔKth and Kfc are attained at hydrogen content in
the metal of 8 ppm. On the whole, for 10Kh15Ni27Ti3W2MoAl steel, at this frequency of
loading and cycle asymmetry, the influence of hydrogen is much weaker than its influence on
the crack resistance under static loading (Figs. 5, 6).

Specific features of the fracture of alloys in hydrogen.

The hydrogenated specimens of 10Kh15Ni27Ti3W2MoAl steel in static tension fail (almost

completely) according to the mechanism of shear accompanied by the formation of a great
number of secondary cracks, in particular, on the grain boundaries (Fig. 7a). In neutral
media, under the conditions of low-cycle bending, we also observe the shear character of the
fracture surfaces with striated structure typical of low-cycle fatigue and the inflections of
striate in passing through the grain boundaries and twins. After hydrogenation, we detect the
intense development of intergranular fracture with intermetallic compounds and carbides on
the fracture surfaces (Fig. 7b). For both loading modes, the difference between the
characters of fracture in the vacuum and in hydrogen is preserved in the entire investigated
temperature range.

In static crack resistance tests, a substantial difference in the character of fracture is

observed in the whole temperature range 293-673 K investigated (Fig. 8a-d). In a neutral
atmosphere, transgranular fracture with a dimple microrelief, precipitations on the bottom of
dimples, and an insignificant number of individual intergranular cracks occur (Fig. 8a, c). In
hydrogenated specimens, in regions near the front of a fatigue crack, facets of intergranular
fracture and regions of quasicleavage with combs of break-off along boundaries carbide-
 7

matrix and intermetallic compound-matrix dominate (Fig. 8b). The combs of break-off
indicate the localization of plastic deformation in thin regions where microcracks are initiated
by quasicleavage on the boundary of the γ and γ΄- phases. In individual regions of fracture
with a plane surface with steps, spalling in a fee lattice was detected. This was caused by
the simultaneous propagation of a crack in several crystallographic planes within the
boundaries of one grain. In the region of shear lips, hydrogen substantially reduces the
fraction and sizes of nonequilibrium dimples and increases the fraction of intergranular
cleavage (Fig. 8c, d).

a b

Fig. 7. Microfractographs of 10Kh15Ni27Ti3W2MoAl steel after tension (a)

and low-cycle bending (b) for the maximum hydrogen degradation (x 1005).

Fig. 8. Fracture surface of 10Х15Н27Т3В2МР

steel at the crack front (a, b) and shearing jaws
(c, d) at static loading at 673 K (a, b) in helium
(a, c) and in hydrogen under the pressure 10
МPа after hydrogenation (623 K; 10 МPа;10 h)
(b, d). × 3000.

Fig. 9. Fracture surface of 10Х15Н27Т3В2МР steel after cyclic loading at 293 K on the deep
( І) (a, b) and upper (ІІІ) (c, d) regions of fatigue crack growth rate diagrams and on
 8

spontaneous fatigue region (e, f) in helium (a, c, e) and in hydrogen under the pressure 10
МPа after hydrogenation (623 K; 10 МPа; 10 h) (b, d, f). × 3000.

Fractures of hydrogenated and non-hydrogenated specimens of 10Kh15Ni27Ti3W2MoAl

steel in the lower portion of the KFFD differ insignificantly. They are characterized by a
specific striated structure with cleavage facets, the number of which is somewhat larger in
the case of hydrogenation (Fig. 9a, b). Moreover, on fractures of hydrogenated specimens,
small numbers of microcleavages, stratifications with slip planes, and twin discontinuities
(tongs) are observed on cleavage facets. In the upper portion of the KFFD, hydrogen favors
the transition from the striated and honeycomb relief of fracture, which are characteristic of
fatigue (Fig. 9c), to the intragranular and intergranular cleavage with plane facets (Fig. 9d).
The largest differences in the character of fracture between hydrogenated and non-
hydrogenated specimens are observed in the stage of final fracture when, in air, dimple
break-off with regions of intergranular fracture (Fig. 9e) and, in hydrogen, a network of
intergranular cracks with individual tracks dominate (Fig. 9f).

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Corresponding author: balitski@ipm.lviv.ua

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