Low temperature and grain size effects on threshold and fatigue crack propagation in a high strength low alloy steel

Materials Science and Engineering, 51 (1981) 203 - 212

203

Temperature and G r a i n S i z e E f f e c t s on Threshold and F a t i g u e C r a c k Propagation in a High Strength Low Alloy Steel Low

J. P. LUCAS and W. W. GERBERICH

Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, MN 55455 (U.S.A.) (Received June 5, 1 9 8 1 ; i n revised form July 8, 1981)

SUMMARY

Threshold and fatigue crack propagation results were observed for a high strength low alloy steel at temperatures ranging from 123 to 300 K and grain sizes ranging from 10 to 123 pro. The threshold stress intensity was observed to increase as the grain size increased. Also a substantial increase in threshold was exhibited as the test temperatures were reduced from 300 to 123 K. The mode o f fracture near and at threshold was ductile for all temperatures and grain sizes. However, at intermediate AK levels and low temperatures the mode o f crack propagation was brittle. Superimposed on the river lines o f the cleaved fracture surface were cyclic cleavage striations indicative of the local fracture mode. The fatigue crack propagation exponent n o f the Paris law increased rather drastically as the temperature decreased and was determined to be directly associated with the brittle fracture process. Emphasis was placed on the subgrain cell structure that develops during cyclic deformation as being a very likely microstructural unit controlling the fatigue threshold process. Thus the grain size influence o f the threshold is partially interpreted in terms o f the subgrain cell size dependence on grain size. The temperature effect on the threshold is discussed from a thermal activation analysis view o f dislocation behavior and plastic flow at low temperatures.

1. INTRODUCTION

The effects of various parameters on fatigue crack propagation (FCP) near the threshold stress intensity level AKTH [1 - 14] are still 0025-5416/81/0000-00001502.50

poorly understood and in some cases unspecified. Considerable effort has been centered around describing the effects of microstructure on fatigue crack growth at threshold. Whereas m a n y microstructural variables have been examined for their influence on threshold, grain size has been given considerable attention most recently [7 - 14]. In general, the threshold is shown to increase, at least for the class of low to medium strength steels, as the grain size increases [9, 10, 11 - 14]. However, there is some controversy over the exact isolated effect of grain size on threshold since, in most alloys, changes in the grain size are usually related to changes in the yield stress [15]. Nevertheless, large grain size effects on threshold are realized. In a ferritic steel, Masounave and Bailon [9, 10] showed that an increase in the grain size of nearly an order of magnitude resulted in an e n h a n c e m e n t of the fatigue threshold by a factor of 2. Priddle [13] investigating a stainless steel also found an increase in threshold with an increase in grain size;virtually no change in threshold was found when the grain size was kept constant while changing the yield stress by 40%. Attempting to isolate the grain size effects, Benson [14] investigated the influence of grain size on threshold in a low alloy steel, while maintaining a nearly constant yield stress, by controlling the precipitation-hardening contribution to the yield stress. Again, the threshold was found to be higher for larger grain sizes [ 14]. Unlike other microstructural and environmental variables, low temperature (i.e. below room temperature) effects of FCP in steels are n o t as prevalent in the literature. This is especially true of studies dealing with low temperature effects on FCP near threshold stress intensity levels. However, low tempera© Elsevier Sequoia/Printed in The Netherlands

204 TABLE 1 The composition (wt.%) of high strength low alloy steel

C

Nb

Mn

S

P

Si

Al

Fe

Grain size (pro)

0.06

0.03

0.35

0.01

0.01

0.03

0.01

Balance

10

ture investigations of FCP above threshold (region II) have been performed on several iron-based alloys [16 - 1 8 ] . In an F e - M o alloy system, Burck and Weertman [16] reported t h a t crack growth rates were reduced as the temperature was lowered from 300 to 77 K. For a series of F e - N i and F e - S i alloys, low temperatures had a very significant effect on the FCP rate. In fact, for these alloys, Gerberich and M o o d y [17] found that the FCP e x p o n e n t n from the Paris relationship [19] da dN

= C AK"

(1)

increased from a b o u t 4 to 20 as the test temperatures were reduced from 298 to 123 K. Excessively high FCP exponents were observed to be associated with the brittle fracture modes which dominate during the fatigue process. Similar results were also observed for a high strength low alloy (HSLA) steel [ 2 0 ] . As acknowledged previously, very few studies exist on low temperature FCP and even fewer results have been reported on the effects o f low temperatures on threshold stress intensity [18, 20 - 2 2 ] . In an HSLA steel, the thresholds increased as the temperature decreased [ 2 0 ] , whereas Stonesifer [18] could find no effect of low temperatures on threshold even at 77 K. Tschegg and Stanzl [ 2 1 ] , however, showed an increase in threshold as the temperature was changed from 300 to 77 K for a low carbon steel. In this investigation, low temperatures ( 3 0 0 - 123 K) and grain size (10 - 123 p m ) were simultaneously examined for an HSLA steel in order to ascertain their effects on threshold and near-threshold stress intensities for constant R ratio and test frequency. Some consideration is given to FCP above threshold at intermediate AK levels at low temperatures and to the predominate m o d e of fracture in this region.

2. MATERIALS AND EXPERIMENTAL PROCEDURES

The material tested in this investigation was an HSLA steel. Specimens were cut from 8.0 mm sheet stock oriented in the longitudinal-transverse position of the rolling direction. The major alloying elements are listed in Table 1, together with the as-received grain size. Grain sizes ranging from 10 to 123 p m were obtained by heat treating high in the austenitizing region at 1473 K for various periods of time in a vacuum of a b o u t 10 -5 Torr. All specimens were then allowed to furnace cool to room temperature. In Table 2 the heat treatment schedule is shown for the various grain sizes achieved. Test temperatures ranging from 300 to 123 K were controlled within an accuracy TABLE 2 Heat treatment and grain size

Gram s ~ e ( p m )

Temperature(K)

Time(h)

10 65 90 123

-1473 1473 1473

-0.5 3 8

to + 1 K by utilizing an insulated chamber. A controlled flow of liquid nitrogen vapor was circulated over the specimen. Also a Nichrome heating element, controlled by a proportional controller in the specimen chamber, assisted in maintaining a constant temperature. The specimen temperature was monitored with a Cu-constantan thermocouple fixed to the specimen. Static testing to determine the tensile properties were carried o u t on an Instron test machine at a strain rate of 5.6 X 10 -4 s-1. Compact tension specimens 7.8 mm thick and surface ground to 600 grit were used for the FCP tests. Fatigue testing was performed on an MTS servohydraulic feedback test machine. A constant:amplitude sinusoidal load was applied to the specimen at a fre-

205

quency of 30 Hz. A fixed load ratio R equal to 0.1 was maintained for all tests. The fatigue crack growth rates were taken at constant Ap over a small incremental crack length Aa determined by compliance techniques [23]. The cyclic stress intensity range was calculated from Y AK = AP ~ (2a)

BW1/2

29.6

3. EXPERIMENTAL RESULTS

3.1. Microstructural and tensile properties

where

Y=

Ap is the applied cyclic load, a is the crack length, B is the specimen thickness and W is the specimen width [23].

\w][a,3/2

( a ) z/2 -

-

185.5|--/

+

+655.7 (2b)

The as-received material consisted of a fine nearly equiaxed ferrite grain structure. Some small colonies of fine-lamellar pearlite could be found scattered throughout the predominately ferritic grains (Fig. I(A)). After heat treating for grain growth, more fine-lamellar pearlite was observed along the grain boundary nodes in the 65, 90, and 123 tzm grain sizes. This is seen in Figs. I(B), I(C) and I(D) respec-

J

(C)

IOOl

Fig. 1. Typical microstructures of the as-received and annealed material showing average grain sizes of (A) 10/,tm, (B) 65 ~tm, (C) 90 p m and (D) 123 ~m.

206

tively. The yield stress is given in Fig. 2 as a function of temperature at various grain sizes. There is little difference in the yield stress versus temperature curves for the 65, 90 and 123 p m grain sizes, and the reason for this is n o t clearly understood. However, the asreceived 10/~m grain size has a much higher yield stress than the coarse grain specimens. This difference in the yield stresses is most probably due to a combination of the precipitation strengthening b y Nb(CN) precipitates and the fine grain size in the 10 p m grain size material.

3.2. Fatigue crack propagation results FCP curves for the 10 ~m grain size are shown at four temperatures in Fig. 3. On examination of the fatigue threshold results at a b o u t 10 -1° m cycle -1 , it is seen that at 300 K the threshold stress intensity AKTH is 8.2 MPa m 1/2. The threshold increased as the test temperature was lowered from 300 to 123 K. At the lowest test temperature the threshold is raised by 5.3 MPa m z/2 at constant grain size. The threshold is enhanced because of low temperature effects by a factor greater than 1.60. The profile of the FCP curves showed a drastic change at intermediate AK levels above threshold, especially at 233, 173 and 123 K. The abrupt change in crack growth rate occurred at a d a / d N value of approximately 10 -s m cycle -1 . Such drastic increases in the FCP rate were related to a change in the fracture m o d e of the propagating crack.

l lO

.~'o~8 .

.

.

.

759

.

The threshold results for the 65 pm grain size can be seen in Fig. 4. Similarly, the threshold increased substantially as the temperature was lowered. At 300 K for the 65 p m grain size, the threshold stress intensity is 9.7 MPa m 1/2. Hence the grain size effect on threshold becomes apparent. There is a 20% increase in

Ld

id 7

E I0

to"

id° 8

10

12 15 20 AK, MPa m~

25

30

Fig. 3. Threshold and F C P curves for the 10/Im grain size samples at 123 K (m), 173 K (D), 233 K (o) and 300 K (e) ( H S L A 2;R = 0.10; frequency, 30 Hz).

d

iv'

100

[]

90

621

580 O 70

483

10' 50

345

40 I

300

510

i 100

150

200

250

500

550

4()0

207

TEMPERATURE, K Fig. 2. Y i e l d stress vs. test temperature for various grain sizes ( H S L A 2): o, 10 pro; o, 65/Jm; A, 90 pro; O , 123 pro.

, ,/,f,lf,,~,l,

10 12

15

,I,L

20

25

Z~, MPo mL~

Fig. 4. Threshold and F C P curves for the 65/.tm grain size samples at 123 K (m), 173 K (D), 233 K (o) and 300 K (o) ( H S L A 2 ; R = 0.10; frequency, 30 Hz).

207

threshold as the grain size is increased from 10 to 65/am. Again, if we start at d a / d N 10 -s m cycle -1 , higher crack growth rates are evident with increasing AK values for the 6 5 / a m grain size. just as was f o u n d for the 10/am grain size in Fig. 3. The FCP curves for the largest grain size, 123 pm, are depicted in Fig. 5. The same trend for d a / d N versus A K is exemplified as before with the 10 and 65/am grain sizes. Once more, a noticeable increase in threshold was witnessed as the grain size increased and as the temperature was decreased. It is worth while to note the combined grain size and temperature effects on the threshold stress intensity. For the grain size-temperature conditions 10/am, 300 K, and 123/am, 123 K, a difference in threshold of 10 MPa m 1/2 was observed. An increase in threshold of this magnitude would n o t be predicted by theoretical models using a dislocation emissary approach [ 24, 25]. Perhaps, the extent of the combined temperature and grain size effect on threshold is better illustrated in Fig. 6 where A K T H is plotted against temperature for the various grain sizes. Fractographs were taken of failed fatigue specimens in an a t t e m p t to discern the type o f fracture involved during the threshold fatigue process. The m o d e of failure at and near threshold was predominantly ductile, although occasional cleaved facets were found

while scanning the breadth of the specimen at threshold [ 20]. The ductile nature at threshold is seen in Fig. 7 for the fine grain size material at 300 K. Surprisingly, the mode of fracture at threshold for a low temperature (123 K) and a coarse grain size (123/am) was also ductile. The fact that we would obtain extensive plasticity at threshold for a coarse grain size material at very low temperatures was n o t anticipated. Figure 8 is typical of the fracture surface observed at threshold at 123 K for the 123/am grain size. The failure appeared to be a ductile transcrystalline type. At intermediate AK levels the shape of the FCP curve change was very obvious especially for 10 and 65/am grain sizes. In this region the FCP e x p o n e n t n increases rapidly as the temperature decreases [ 17, 26 ]. This acceleration in growth rate is emphasized in Fig. 9. 20

i

i

i

,

~8 16

10 B I

o

50

i

Ioo

400

"I'EMPERATURE, K

,d

Fig. 6. Threshold vs. t e m p e r a t u r e for 10 # m (o), 6 5 / a m ([]) and 1 2 3 / a m (o) grain sizes ( H S L A 2; R = 0.10; f r e q u e n c y , 30 Hz). iO-7

~ l O -a

id a

io'e . . 10 12

15

20

25 "N:?

Z~K, M ~ m ~'~

Fig. 5. Threshold and FCP curves for the 1 2 3 / a m grain size samples at 123 K (1), 173 K (D), 233 K (o) and 300 K (o) ( H S L A 2 ; R -- 0.10; f r e q u e n c y , 30 Hz).

Fig. 7. A fractograph of fine grain (10 Pro) material observed at threshold at 300 K.

208

Fig. 8. A fractograph of coarse grain ( 123/~m) material observed at threshold at 123 K.

Fig. 10. A fractograph observed in region II (AK = 20 MPa m 1/2) for the 123 pm grain size at 123 K, exhibiting cyclic cleavage markings along river lines.

. x t o "~

cyclic cleavage markings are perpendicular [17, 26] with respect to the characteristic river lines indigenous to brittle fracture failures. However, it should also be n o t e d that the microscopic crack growth rate due to cyclic cleavage processes is much higher than the macroscopic crack growth rate [ 2 6 ] . In Fig. 10, forexample, da/dNimicro ~ 10 -e m cycle -1 while da/dNl~acro ~ 10 -s m cycle -1.

E Z

i x l O"e

4. DISCUSSION

16

18

2

24

Io

J is

I

I I

m 22 ~4

|

I

18

~e

';0 24

AK, rvPom i/2 (a)

(b)

(c)

Fig. 9. A plot of da/dN against AK emphasizing crack growth in region II and the variation in the FCP exponent n with temperature for (a) the 10/~m, (b) the 65 pm and (c) the 123 jura grain size samples: m, 123 K; [], 173 K; o, 233 K; $, 300 K. Here the FCP e x p o n e n t for the coarse grain size was 7.9 at 300 K and a b o u t 19.6 at 123 K. The accelerated crack growth in this region is associated with the onset of extensive brittle fracture [ 17, 20, 2 6 ] . Transgranular cleavage fracture was observed as the temperature was reduced. However, closer examination of fractographs revealed details associated with cyclic cleavage markings. Cyclic cleavage markings are clearly exhibited in Fig. 10. The

4.1. Grain size effect o n AKTH By itself, the grain size affects the observed threshold stress intensity of HSLA steel. In order to elucidate the isolated grain size dependence on threshold, AKTH was plotted as a function of the yield stress for t w o grain sizes in Fig. 11. For the 65 and 123 p m grain sizes the yield stress versus temperature curves (Fig. 2) were virtually identical, which is implicit in Fig. 11 because the 65 and 123 # m grain size curves are almost parallel. Consequently, the isolated grain size effect on threshold is evident. In a previous investigation, Masounave and Bailoh [9, 10] found a similar effect of grain size on threshold in a ferritic steel. They showed that, although the yield stress remained nearly constant for coarse grain sizes, AKTH increased as the grain size was increased by a factor of a b o u t 2. They rationalized that the grain size effect on threshold was due to the deviation in the crack path as it propagated through grains of

209 (~APa) 276 ,

414 ,

552 ,

,

690 ~

i

828

I00

I10

20 18 16

E

B 6 3,O

40

50

6

Fig. 11. Threshold

80 0 Tlmp, ISTRESSlo.2~, K$i

vs.

120

temperature-dependent yield

stress for the 65 pm (A) and 123 pm ($) grain sizes (HSLA2). different sizes. According to Masounave and Bailon [9, 10], a compatibility between the propagating crack front and certain crystallographic slip systems must be satisfied for crack growth; the probability that the necessary conditions are fulfilled increased as the grain size decreased. However, although such an approach seems reasonable, it is probably more applicable to the fatigue crack growth rates than to the threshold [14]. In order to account for the grain size effect on threshold in a stainless steel, Priddle [13] considered the crack path morphology and the macroscopic growth rate near threshold. He concluded that crack path deviation from the plane of maximum tensile stress led to a reduction in the effective stress intensity to a certain degree [13]. An additional reduction in the value of AK was thought to be associated with the increase in real crack surface area compared with the observed macroscopic crack length and specimen width. Although these factors no doubt contributed to observed results, they could not account for the observed magnitude of grain size on threshold in some low alloy steels [9, 10, 12 - 14, 20]. Gerberich and Moody [22] modeled the threshold stress condition by considering that a semicohesive zone, which consisted of ligaments commensurate with the grain size, developed because of irregularities in an equilibrium crack front. Thus the ligaments acted as tractions to reduce the stress intensity at the crack tip, thereby increasing AKTH.

Such a model predicted AKTH reasonably well for a number of alloys. Recently Ritchie [11] suggested that the threshold is controlled by oxide-induced crack closure, especially at low R ratios. With such an approach the grain size effects appear to be minimal. However, grain size effects are still observed at low temperatures [20], where the oxide effect on closure would be much less because of the gettering of oxygen on the walls of the cold chamber. The previous models [9, 13, 22] describing the influence of grain size on fatigue threshold stress intensity are all based on a common feature. They all rely on the concept that the stress intensity at the crack tip is effectively reduced either by crack path deviation [ 9, 13] closure [ 11] or by ligaments left by nonuniformity in an advancing crack front [22], and hence an increased AKTH.However, the approach to describing the grain size effect on fatigue threshold in the present investigation is quite different. The basis of the present model is primarily concerned with the subgrain cell structure that develops within the grains in the vicinity of the crack tip as a result of the applied cyclic load. Iron and low alloy steels in general show a high propensity for subgrain cell formation during cyclic plastic deformation [27 - 29]. To rationalize the effect of grains that are considerably larger than the cyclic plastic zone on threshold, it is presumed that the subgrain cell structure developed during cyclic deformation is the controlling microstructural unit. The fact that the subgrain might influence the threshold has also been suggested by McKittrick et al. [30]. It is also presumed that the subcell structure that forms is influenced by the initial grain size in the material. Indirect evidence of a grain size-subcell dependence has been deduced from the grain size effect on the cyclic strain-hardening exponent tic [31]. Realizing that geometrically necessary dislocations emitted by grain boundary sources depend on the grain size [32], we can use the following approach to describe the j3c-grain size dependence. At saturation for a constant cyclic plastic strain amplitude, a likely configuration would be one in which the dislocations are arranged in stabilized subgrain cells spanning the entire grain. Additionally, if the dislocations spacings within the stabilized subgrain cell structure were nearly equal, then the size of the subgrain cells that developed

210 during cyclic plastic deformation can be shown to be a function of the initial grain size of the material. The fact that the subgrain cell size L is dependent on the grain size D perhaps can be realized simply by noting the influence of L and D on the geometrically necessary dislocation density pG during plastic deformation. For an inhomogeneous (polycrystalline) material plastically deformed to a certain strain [ 32 ], pG ~

4Me

crock

(3)

bD

where M, e and b are Taylor's orientation factor, the true strain and Burger's vector respectively. Also, it can be shown that, for a two-dimensional square net of subgrain cells [33], 2 pG ~ __ Ll

(4)

where I is the dislocation spacing in the subgrain cell wall. Now, if we combine the two preceding expressions for pG, it is resolved that the subgrain cell size is given by b L ~ 2Mel D

(5)

as a first approximation. (It should be pointed out that l is a function of the material stress and hence also a function of the grain size. Nevertheless, the net result is that the subceU size will still increase with increased grain size as will be discussed in a future paper [34] .) Hence the subgrain cell size at saturation would be larger for material of larger grain size deformed cyclically at a constant strain amplitude. Thus for a fine grain material the volume of strain-hardenable material available for subsequent hardening at higher strain amplitudes would be less than the volume of strain-hardenable material for the coarse grain size specimen. This then is a possible explanation of the/~e-grain size dependence [31]. It is envisaged that, at the threshold stress intensity level after numerous fatigue cycles, the subgrain cells achieve a critical (saturation) size with impenetrable cell walls. This is depicted in Fig. 12. At threshold, dislocations being emitted from the crack tip or dislocations at a distance for the crack tip are simply shuttled back and forth within the boundary of the cell wall in compliance with the frequency of the applied load. In other words,

e

Fig. 12. A schematic depiction of the substructure that develops at threshold during cyclic deformation. the cell walls act as formidable barriers to dislocations approaching the cell wall as well as to dislocations attempting to move within the walls. It should be noted that, as the saturated cell size becomes larger, the slip distance can accommodate additional dislocations. Consequently, the back stress would be higher (and hence a higher AKTH ) because of an increase in the n u m b e r of dislocations in slip band pile-up [35, 36]. The threshold condition will persist until the applied stress is increased to a level sufficient to break down the stable subgrain cell structure. When an ample stress intensity has been attained, crack propagation will proceed and perhaps some other microstructural feature will play a controlling role. 4.2. L o w temperature effect o n AKTH The strong temperature dependence of AKTHis exemplified in Fig. 6 for several grain sizes. Obviously, the change in elastic modulus with temperature cannot account for the observed increase in threshold [25]. The temperature dependence on threshold can be deduced from a thermal activation analysis for plastic flow. The stress required to produce plastic flow (slip) becomes greater as the temperature is reduced below 300 K in low alloy steel. This is because thermal energy is lost at low temperatures which can assist the m o v e m e n t of dislocation over short-range or thermal barriers [37, 38]. These short-range or thermal barriers that mobile dislocations must overcome are usually considered to be

211 the Peierls-Nabarro stress associated with intrinsic lattice friction, impurity atoms and dislocation nodes. These barriers manifest themselves at low temperatures b y an increase in the thermal c o m p o n e n t of the flow stress. The flow stress can be expressed in terms of t w o components: anow(T) = o * ( T , ~ ) + ai(ep)

(6)

where o* is the thermal or effective stress and ai is the athermal or long-range stress. Since the reversed plastic zone size is inversely proportional to the square of the flow stress, AKTH(T ) should be influenced directly by A o * ( T ) provided that the mode of fracture at threshold is the same at all temperatures. Hence AKTH increases as AaTH(T ) increases at low temperatures, specifically because of the increase in Aa*(T). 4.3. Fatigue crack p r o p a g a t i o n at i n t e r m e d i a t e A K and low temperatures

The high FCP rate and the corresponding high FCP e x p o n e n t n is due to the occurrence of a brittle fracture process at high AK values at low temperatures [17, 20, 26]. The mode of fracture at intermediate AK levels and lower temperatures has been identified as cyclic cleavage fatigue crack growth [ 2 6 ] . It is shown in this investigation and elsewhere [20, 26] that da/dN[micro for cyclic cleavage crack growth rate is as much as two orders of magnitude higher than da/dN]macro for the macroscopic crack growth rate. The crack growth rate can be described by a proposed model [ 26] where cyclic cleavage is interpreted in terms of the local crack-tip-opening displacement (CTOD). This is CTOD controlled by the necessary dislocation density required to produce river lines of a certain height and the average velocity of dislocations emitted to some distance from the crack tip. Such a model was able to predict the increased crack growth rate at low temperatures as the transition in fracture modes went from a ductiledominant to a cleavage-dominant process for FCP. Obviously , the incorporation of such features as cyclic cleavage, cyclic plasticity, grain size and temperature into a single model for predicting FCP at low and intermediate AK levels is well b e y o n d our present understanding.

5. SUMMARIZING REMARKS AND CONCLUSIONS (1) The threshold stress intensity increased as the grain size increased from 10 to 123 ~m in an HSLA steel. (2) The threshold was observed to increase as the test temperatures were reduced from 300 to 123 K because of the increase in the thermal stress c o m p o n e n t of the flow stress with decreasing temperature. (3) The crack growth rates and the FCP e x p o n e n t n were higher because of the cleavage-dominated fracture mode at low temperatures and intermediate AK levels. (4) Although previous models suggest that the threshold is controlled by mechanisms such as crack path deviation, traction ligaments due to crack path tortuosity or oxide-induced crack closure, an additional factor is presented here. This describes the grain size effect in terms of the subgrain cell structure that develops in the material during cyclic plastic deformation and is believed to be a contributing factor also. ACKNOWLEDGMENTS Financial support for this research was provided by Grant DOE-DE-AC02-79ER10433 from the U.S. Department of Energy. Thanks are also due to Dr. P. Mangonon, formerly of Inland Steel, and Inland Steel for supplying the material used in this investigation. REFERENCES 1 R. J. Bucci, W. G. Clark, E. T. Wessel and T. R.

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