Dynamic recovery and strain hardening in the hot deformation of type 317 stainless steel

Materials Science and Engineering, 81 (1986) 259-272

259

Dynamic Recovery and Strain Hardening in the Hot Deformation of Type 317 Stainless Steel N. D. RYAN and H. J. McQUEEN

Mechanical Engineering Department, Concordia University, Montreal, Quebec H3G 1M8 (Canada) E. EVANGELISTA

Dipartimento Scienze Materiali, Universit~ di Ancona, Ancona 60128 (Italy) (Received January 2, 1986)

ABSTRACT

Worked specimens o f type 31 7 stainless steel, deformed in torsion to fracture in the temperature range 9 0 0 - 1 2 0 0 °C and at strain rates o f O. 1-5 s -1, exhibited flow curves which increased to peaks and decreased to steady state regimes. As a result o f dynamic recovery, the flow stress depended on the strain rate through the hyperbolic sine function and on the temperature through an Arrhenius function with an activation energy o f 496 k J mo1-1. The strain-hardening rate 0 was analysed as a function o f stress a to derive the saturation stress as* due to dynamic recovery alone. Extrapolation o f the O-a curves to a = 0 determined the athermal hardening rate; extrapolation o f the o * - T curves to T = O and o f the o~*-~ curves to em determined os* at zero recovery where a is insensitive to strain rate and temperature respectively. The subgrain sizes in the steady state regime were related by simple functions to the T and ~ conditions, to the flow stress and also to the room temperature hardness. The behaviour o f type 3 1 7 steel is compared with that o f type 304 steel which has been the subject o f much research.

1. INTRODUCTION

The behaviours during the hot working of austenitic stainless steels are of interest beyond their industrial significance because they are f.c.c, iron alloys which do not transform on cooling. It is expected that they have considerable similarity in substructure to that of carbon steels even though they have a somewhat decreased stacking fault energy [1-4]. More0025-5416/86/$3.50

over, the high solute levels retard static recrystaUization, making it easier to study both dynamic recovery and rebrystallization. These mechanistic features cause stainless steels to be much stronger and less ductile at elevated temperatures than are simple carbon steels. The hot-working behaviour of type 304 stainless steel has been subject to some dozen fairly thorough studies [1, 5]. However, type 317 stainless steel has been given only two or three such examinations [1, 5-10]. In this steel the high molybdenum content leads to solute and precipitation effects which make it considerably more sluggish in dynamically recovering and recrystallizing than type 304 steel and consequently endow it with a higher hot strength and a lower hot ductility, thereby giving rise to processing problems. The austenitic stainless steels are ideal materials for the study of hot-working mechanisms and the influence on them of solute and precipitates [1, 2, 4, 11-14]. In this paper, part of a project to s t u d y types 301, 304, 316 and 317 steels in both the as-cast and worked conditions is reported. The mechanical behaviour of as-cast type 317 steel has been described in previous reports [5, 6], and so the present paper will concentrate on the worked material. Moreover, some comparisons will be made with type 304 stainless steel as a model material. In light of the theme of the conference published in this volume, the integrative concept is the influence of dynamic recovery on the elevated temperature mechanical properties. One property which is not considered is the ductility since it is affected principally by the dynamic recrystallization and only secondarily by recovery [12]. © Elsevier Sequoia/Printed in The Netherlands

260 I

2, E X P E R I M E N T A L T E C H N I Q U E S

I

I

I

|

"C

Type 317 steel provided b y Atlas Steels of Tracy, Quebec, had a composition of 18.60 wt.% Cr, 13.88 wt.% Ni, 3.22 wt.% Mo, 1.73 wt.% Mn, 0.44 wt.% Si (the sum of the metallic solute content is 37.87 wt.%), 0.035 wt.% C, 0.017 wt.% N and the balance iron. The specimens had gauge dimensions of 25.4 mm length and 6.25 mm diameter. They were annealed at 1050 °C for 30 min to develop grains of 57 #m (as cast) and 72 #m (worked) and were quenched to avoid carbide precipitation. Torsion tests, commencing with preheating for 5 min in a radiant furnace, were conducted in a closed-loop servo-controlled hydraulic hot-torsion machine [5, 6]. The following relationships were used to calculate the surface shear strain 7 and the average shear stress r from the torque F' [5, 6] : 7 =

21rr

I

revolutions

= 31/2e

(la)

m

r = 2~rS (3 + n" + m) o

31/2

(lb)

The value of the strain-hardening coefficient n f! was approximated as zero near the peak stress and the strain rate sensitivity m was evaluated across the testing spectrum (Fig. 1) [8]. The equivalent stress and strain were calculated from the von Mises criterion (eqn. (1)).

Specimens for optical microscopy were obtained from longitudinal, transverse and tangential sections. The last provides the clearest indication of elongated grains. Observations in the first two are restricted to the region just below the surface where ~ is almost the same as the values quoted. The grain sizes were determined b y the three-circle intercept m e t h o d in A S T M Standard E 112 [15]. The slices for transmission electron microscopy (TEM) specimens were cut parallel to the axis by means of a slow speed diamond saw. They were ground down to a uniform thickness and then chemically thinned. The final thinning was mainly b y double-jet thinning of disc specimens. The examinations were conducted at

!

n 14'

12"

900

92

~"Q"''~

I0"

I m 'OST '109 "091

806O"

~

'

t

5

6

~4ot,- 3,0-

2D-

15-

.2

.~d "~"

.I I 0.1

I I I 0.360.5 I STRAIN RATE ~, s -I

I I ~635

I I0

Fig. 1. The strain rate sensitivity m of the torque appears as the slope in this log r vs. log e plot for as-cast (m, ) and worked (n, - - -) type 317 steel. r

=

A~ m

=

BF n

As t h e t e m p e r a t u r e increases, 1~ decreases b u t m increases. T h e stress e x p o n e n t n (eqn. (2)) is t h e reciprocal o f m. T h e results o f R a d u e t al. [8 ] (~, ) are for a c o m p a r a b l e steel.

200 kV. The subgrain sizes were determined from a linear intercept analysis normal to the long axis. These measurements were multiplied b y 1.68 to make them comparable with results from areal analysis [16J.

3. R E S U L T S

The torque-strain rate curves portrayed in Fig. I determine the value of m for the calculation of stress which compensates for the strain rate gradient. The strain rate sensitivity m increases with increase in temperature. Data from Radu et al. [8] for a steel containing 17.28 wt.% Cr, 12.85 wt.% Ni, 3.66 wt.% Mo (the sum o f the metallic solute content is 35.90 wt.%) and 0.03 wt.% C are also presented; b o t h the torque levels and the strain rate sensitivities are in reasonable agreement. The equivalent stress-equivalent strain curves ( o - e ) appear in Fig. 2. In addition, data for t y p e 304 steel (containing 17.6 wt.% Cr, 0.27 wt.% Mo, 8.52 wt.% Ni ( the sum of the metallic solute content is 29.16 wt.%) and 0.07 wt.% C) [3] is included showing that it is not as strong as t y p e 317 steel. From the flow curves for t y p e 317 steel, 0-a plots are de-

261 450

20C

~P 360 :E

18C

uJ 16C

~

270

14C

tBC

6

~

12C

~

100

& ~ 9o

I ._.,o=o

t--

40 60 80

so

......

:.I

.J

~, sc 50 "' 4C

450~L-/~-"

120 160 ZOO 240 FLOW STRESS~p,MPa

Ca]

" ~ I~,OO o ?

36o.

~oo-c-

'9OO'C-~I !I" ' ' 'ic ~"

Hill

I

Z80

:32.0

'

1

I

i

20 27C -

i

i

20 EQUIVALENT STRAIN

Fig. 2. Representative equivalent stress-equivalent strain curves for type 317 steel as-cast ( , - - -), for type 317 steel from ref. 10 ( ) and for the softer type 304 steel from ref. 3 (--.--)where ~p is the strain to peak, e i the fracture strain and ec the critical strain for dynamic recrystallization: - - , , , steady state stress os; - - -, saturation stress os*. The decrease in strain hardening with increase in T or decrease in ~ is due to dynamic recovery, which would lead to the curve saturation at os* in the absence of dynamic recrystallization, giving rise to the peak and flow softening.

r i v e d w h e r e 8 (= da/de) is t h e r a t e o f s t r a i n h a r d e n i n g (Fig. 3) [ 1 7 , 1 8 ] . T h e s e s h o w a rapid d e c r e a s e in 0 w h i c h is a c c e n t u a t e d b y a h i g h e r t e m p e r a t u r e T o r a l o w e r s t r a i n r a t e %. E x t r a p o l a t i o n o f t h e 0 - 0 c u r v e s t o 0 = 0 det e r m i n e s a s a t u r a t i o n stress o,*. H o w e v e r , 0 actually decreases more rapidly because of dynamic recrystallization to reach 0 = 0 at or, t h e p e a k stress o n t h e o - e curves. F r o m t h e 0-0 extrapolations, stress-strain curves result. ing e n t i r e l y f r o m d y n a m i c r e c o v e r y c a n b e d r a w n (Fig. 2). The p e a k stress ap (corresponding to ep in

Fig. 2) is dependent on temperature T and strain rate ~. Several mathematical functions have been used for this in the literature [4-6].

=°,8c / f l l T I , ~F' II II I! ET ;T AL. At,

0,

40

!D)

~, %

8o

120 leo 2oo FLOW STRESS /~,MPo

240

28o

Fig. 3. Curves of strain-harder~ng rate 0 us. a for (a) as-cast and (b) worked type 317 steel, indicating the critical stress o e for dynamic recrystallization, the peak stress op (experimental, 8 = 0) and the saturation stress os~ (extrapolation to 8 = 0): , , @= 0.1 s-l;- - - , @ = 1.0 s - l ; , ~ = 5.0 s -1.

where 80 = 4586 M P a for as-cast type 317 steel, 80 ~ 4295 M P a for worked type 317~steel and 80 3395 M P a for type 804 steel.As Z decreases, increased d y n a m i c r e c o v e r y l e a d s t o a m o r e r a p i d d e c r e a s e in 8

and a lower value of Gs*. Additional data for type 304 steel (o) exhibit a higher level of dynamic recovery [17]. In the data of Cars et al. [18] the change in slope indicates the start of subgrain formation. F o r s o > 1.2, ~ = .4' e x p ( ~ o ) e x p ( - - R - ~ )

(3)

Also

~ = A " sinh(ao)"'exp(--~--J

(4)

and

For ~ < 0.8,

= Aan e x p ( - - ~ )

(2) = f.(o)

(5)

262

where A, A ', A ", ~, [3, n, n' and Q are empirical constants and Z is the Zener-Holloman parameter. The traditional creep p o w e r law is illustrated in Fig. 1 where m = 1In [5]. It is seen that n comes close to the classic creep values of 4 - 5 at 1200 °C (n = 5.3) b u t increases to 17 at 900 °C. The exponential law appears in Fig. 4 and was effective for the stronger cast materials [6], b u t for the worked stainless steels it was unsatisfactory in the low 6, Z and o corner [5]. For as-cast t y p e 317 steel and worked t y p e 317 steel, Q is 508 kJ mo1-1 and 496 kJ mo1-1 respectively, which is much higher than for molybdenum-free t y p e 304 steel where Q was about 400 kJ mo1-1 [5, 6]. The hyperbolic sine law is illustrated in Fig. 5 which also shows the determ i n a t i o n o f the activation energy. Additional data obtained by Dhosi et al. [9] for a steel

1

I

IlOOeC

1

IO00*C

I I I

900°C

II

I I I

I

II

I

~/, l U l l / I i~1 l i ~_

I,

I! l ! l I// l /~'I

/ / /'/ /// / /r/ /,',:'1

lg,// I /,!'l /

/

'

/ ,

l I

I

I

!O0"C I

60

I

a

I

I

I

I

180 300 420 EQUIVALENT FLOW STRESS ~'(MPo)

Fig. 4. I n a p l o t o f log ~ vs. o, t h e lines at each T for t h e cast alloys are s t r a i g h t a n d parallel, c o n f i r m i n g eqn. (3). F o r t h e as-worked alloys t h e h i g h e r d y n a m i c r e c o v e r y a t low Z gives rise t o a c u r v a t u r e . T h e cast m a t e r i a l e x h i b i t s a higher s t r e n g t h b e c a u s e 5 p h a s e particles r e t a r d recovery. C o m p a r a t i v e d a t a f o r t y p e 304 a n d 3 1 6 steels [5, 6] are s h o w n . Symbol

A s cast or workea

Steel

o, - A, - A, B,--~, - - -

As cast As cast Worked As cast Worked

Type Type Type Type Type

~ (MPa -1) 304 316 316 317 317

5.4 × 10 -2 5.4 × 10 -2 4 . 8 × 1 0 -2

containing 16.7 wt.% Cr, 15.0 wt.% Ni, 4.06 wt.% Mo (the sum of the metallic solute content is 37.73 wt.%) and 0.050 wt.% C deformed in tension is included. The decrease in the yield stress a r and the steady state stress Os with decrease in Z is shown in Fig. 6 for as-cast type 317, 304 and 316 steels to clarify the increase in the amount o f strain hardening [6]. It can be seen that, in the warm range, o r is almost athermal whereas the steady state stress as decreases rapidly, relative to oy. The decrease in the gap a m - - oy between them, which continues into the hotworking range, is indicative of the rapid increase in dynamic recovery; dynamic recrystallization also makes an additional contribution. The T and ~ dependences of the saturation stress are shown in Fig. 7 and Fig. 8 respectively. The logarithm of the saturation stress is plotted against T to yield straight lines, one for each rate, which converge to a point os0*. In contrast, am* is related to ~ by the p o w e r law. The values of n vary from 5.3 to 17 and the constant-T lines converge to a point at era = 1.4 X 10 ~ s-1 and asm* = 520 MPa. Optical microscopy of the specimens deformed b e y o n d the peak revealed that dynamic recrystaUization had taken place (Fig. 9). The relatively limited strain (e~ ~ 1.6) was able to break up the cast structure and the solution of much of the ~ phase. The grains were almost equiaxed a n d w e r e smaller than the initial grain size. The dynamic recrystallization grain size is inversely related to t h e conditions of deformation as given b y Z. The electron micrographs (Figs. 10 and 11) reveal subgrains that range from elongated at high Z to equiaxed at low Z. Across the entire Z range at 1000 and 1100 °C, the walls are narrow, showing some neat arrays and the cell interiors have low dislocation densities. Within each specimen, identifiable grains exhibit a broad variation in substructure shape and orientation. The elongated cells are more frequent at high Z (Fig. 10(a)) b u t become divided b y regularly spaced transverse boundaries which lead to cubical subgrains at medium Z (Fig. 10(b)). At low Z, equiaxed subgrains are much more c o m m o n and make it more difficult to distinguish the dynamic recrystallization grains. In a few cases, recrystallized nuclei are present; the small nuclei, only slightly larger than the subgrains, could have

263 i

'

I'

'

~

'

Symbol 4111 1.0- I 0 .69- "~0 ~, .'S5 . 3 . 6 ,25- ~.1.8

Steel

Reference

As cast

Type 317 Type 317 Type 317 Fe-16.7wt.%Cr15.0wt.%Ni4.06wt.%Mo-

This study 4.0 This study 4.5 4.5 [8] 4.5 [9]

--[>-- Worked Worked As cast

~ o. ~Lo ~ - .44, -I,0

As cast or worked

0.1

laoo

-4

-~, 5'o

25

~a)

,oo

,o0o

6

.'2 ,~o

Worked

~o ,c

.~,

.'6

LOG Lsinh rj ~'p)

,s'o

~o

i

18

iO

BOUIVALENT FLOW STRESS O'o,MPa

2~o

112

th4

3Go

320

0.050wt.%C Fe-15.8wt.%Cr14.0wt.%Ni4.3wt.%Mo0.027wt.%C0.14wt.%N

[10]

l.d

300. 1.1

@P 250- L(

e

Ol

~200ne I-¢0

o.q

150-0.4 Lu I-Z

I00-

.J ~

o.=

As cast or worked

Steel

Reference

---m--

As cast

--o--

Worked Worked

Type 317 Type 317 Type 317 Fe-16.7wt.%Cr-

This study 508 This study 496 [8] 503 [9] 502

Symbol

0

5 0 - o.~

OA

As cast

25-

15.0wt.%Ni4.06wt.%Mo0.050wt.%C

O.tl

0.66

OTO

(b)

074. (:178 O82 V'TxlO 3, K-Z

0.86

0.90

QHW (kJ mo1-1)

0.94

Fig. 5. (a) A plot of log e vs. log{sinh(c~a)} in which the lines for both as-cast and worked type 317 steel at constant T are parallel and straight in confirmation of eqn. (4). e = A(T) {sinh(~p)} n where ~ = 1.2 X 10-2MPa-1 and n is the stress exponent. (b) A graph of log {sinh(~a)} vs. l/T, in which the lines for constant ~ are straight and parallel.

(

= A~sinh(~p)} n exp ~-- - - ~ - ]

~ = A(~) + QHWRT

giving the activation energy QHW as 2.3nR (the slope). The slopes of these lines show that Q for the as-cast steel is higher than that for the worked steel, indicating reduced dynamic recovery. For comparison the data for two similar steels [8, 9] have been added. f o r m e d d y n a m i c a l l y a l t h o u g h lacking subs t r u c t u r e (Fig. l l ( a ) ) . Large nuclei h a v e clearly g r o w n a f t e r d e f o r m a t i o n ceased, having f e w d i s l o c a t i o n s a n d s h o r t s h a r p s t r a i g h t t w i n s (Fig. l l ( b ) ) . In o n e s p e c i m e n a region consists e n t i r e l y o f s t a t i c a l l y r e c r y s t a l l i z e d

grains a s s o c i a t e d w i t h stringers o f 8 phase, w h i c h c o n t a i n a d y n a m i c r e c o v e r y substructure. T h e l o g a r i t h m o f t h e subgrain size ds is p l o t t e d against t h e l o g a r i t h m o f t h e s t e a d y s t a t e f l o w stress in Fig. 12. T h e slope is a b o u t

264

300

~ |

30C .2.5 25C

u~

(h

~5c ,W E

"C 900

-°

0

9 n~

1

-

12 0.083 -D 0090 I

1

~

0

-''-~

~11

I100

"" i

j

7.3 0133

,oc • 2.G -~.e

t2 o o

6,5 0.154-D

5c g 53 0189 o 3C

~

1.5

I00

I

o'.,

(a)

STRAIN RATE ~,s "z

5.0

IOOO 8 0 ~ . , 9 ~ 0 I 0 0 0 I1001200 T i C for 1.0 se¢-I

304,316

~oo ,~o ~ o ,,~o~ o 317

1400

6OO

TIC 'for 1.0 OIll:-I

400

TEMPERATURE-CORRECTED STRAIN RATE Z, lSC"~

Fig. 6. The variation in the yield stress Gy and the steady state stress as for as-cast type 317 stainless steel (o) with Z through the use of the activation energy from Fig. 5(b). The narrowing of the gap O s Oy as T increases and e decreases is principally the result of dynamic recovery. Data for as-cast type 304 steel (o) and type 316 steel (~,/x) are added for comparison [ 6 ]. -

-

~,300 u~250

~2oo 150

-$jOO ~ 5o 30 (b)

i

I

i

i

" o~i lil-g

~sOO

.bo

i

i

i

i

I i 0 -I

I i0 0

I I I i01 I0 2 I03 STRAIN RATE I~, s-I

I i0 4

I 105

Fig. 8. (a) The plot of log a* us. log ~ giving straight lines for each T with the slope m increasing as T increases for as-cast type 317 steel (m) and worked type 317 steel ([]) (this is similar to the breakdown in the power law relationship in Fig. l(h); (b) however, according to eqn. (8) the lines converge at em where stress is independent o f T.

~* =A~m

where m (= l / n ) is the strain rate sensitivity and n is the stress exponent.

0

I

I

273

473

I 673

I

I

873

1073

TEMPERATURE,

I

I

I

1273

t473

1673

K

Fig. 7. The plot of log os* vs. temperature provides straight lines converging to a point aso* at 0 K according to eqn. (8). The data for as-cast type 317 steel (ai,-....) are slightly higher than the data for worked type 317 steel (B, ). Comparative data [17] for type 301 steel (<), ) and type 304 steel (o, - - -) are provided. The saturation stresses oso* at 0 K for type 317 steel, type 301 steel and type 304 steel are 28.2 × 10 s MPa, 14.1 × i0 ~ M P a and 7.3 × 103 M P a respectively.

-- 1 in agreement with previous results for type 304 steel [16, 1 9 - 2 1 ] . In addition, the reciprocal of the subgrain size is proportional to log Z as has been observed in type 304 stainless steel (Fig. 13) [ 1 9 - 2 3 ] . Because of the differences between the activation energies for type 317 and type 304 steels, their Z values are quite different and lead to a separation o f data which makes comparison difficult as in Fig. 6. It was thus decided to shift the Z scales to bring them into coincidence at a specific working condition: 900 °C and I s -z.

265

the subgrain size increases with increasing T or decreasing 6. In comparison with the subgrain sizes reported for type 304 stainless steel [19-23], the subgrain sizes for type 317 steel are slightly smaller but have the same dependence on Z. The relationship of the subgrain size to the high temperature flow stress is the fairly traditional one: os ~ ds -1

Fig. 9. Optical micrographs o f (a) t h e initial dendritic structure and (b) the final d y n a m i c recrystallization grains with ~ stringers at a fracture strain o f 1.6 for 1000 oC and 1.0 s- 1 . (Magnifications: (a) 4 0 × ; ( b )

80x.)

This had been selected previously for the comparison of many steels because there the properties differ more than at higher temperatures [5, 24]. The hardness of the product at room temperature is related to the hot-worked subgrain size according to the formula H = H o + k d s -1"5

(6)

where Ho is the annealed hardness and k is a material constant.

4. D I S C U S S I O N

The increase in dynamic recovery with decrease in Z is clearly demonstrated by TEM which shows the progress to a larger and more polygonized substructure. The quantitative expression of this appears in Fig. 13 where

(7)

From Fig. 12, it is seen that for a given sui~ grain size the flow stress in type 317 steel is greater than in type 304 steel: Thus, it appears that the strengthening arises in part from the direct action of solutes and precipitates and in part through the fact that solutes and precipitates alter the substructure to be smaller at a given deformation condition. Thus, from Fig. 12 at 900°C and 1.0 ~-1, the subgrain sizes ds of type 317 steel and type 304 steel are 1.11 #m and 1.21/~m, which give strengths of 275 MPa and 195 MPa respectively. The retained substructure gives rise to increased hardness at room temperature (eqn. 6). The power of ds is 1.5, higher than the values of unity commonly assigned to aluminium and ~-Fe alloys [12, 16, 26, 27] ; however, it has been observed in Al-Mg alloys [28]. This differs from the traditional HallPetch power of 0.5, because of the increase in strength of the subboundaries as the cells become smaller [12, 16, 27]. In addition, since this material has undergone dynamic recrystallization, the grain size decreases as the subgrain size decreases and thus contributes to the increased effect of subgrain size. The dynamic recovery is also evinced by the flow curves and the strain-hardening rates. Inspection of the curves in Fig. 2 shows that the rate of initial hardening and the maximum level of the flow stress decreases as T increases and ~ decreases. This is evident in the peak stress op, the steady state stress os and the saturation stress a~*. Although dynamic recrystallization is responsible for the characteristic shape of the flow curves, it contributes only a small amount of the overall softening when the differences o*(T) -- as(T) (Fig. 2) and a*(20 ° C ) - a * ( T ) (Fig. 7) are considered. The rate of strain hardening f o r t y p e 317 steel is greater than for type 304 steel (Fig. 2, the curves from McQueen e t al. [3] are those of the TEM specimens of Fritzmeier e t al. [19] in Figs. 12 and 13). For both as-cast and

266

Fig. 10. TEM substructures for type 317 steel, illustrating the decrease in dimensions and increase in aspect ratio as Z increases through the range (a) 1100 °C and 0.I s-l, (b) II00 °C and 1.0 s-1, (c) ii00 °C and 5.0 s-I and (d) 1000 °C and 1.0 s-I.

worked type 317 steel, the strain-hardening rate is higher than for type 304 or type 316 steel with the result that the peak strain, indicating the onset of dynamic recrystallization, is lower [5, 6]. Part of the increase in strain

hardening, notably in the worked steel, is related to the higher solute content. However, the combination of a higher ap and a lower ep is due to the presence of 5 particles (many in the as-cast and very few in the worked

267

Fig. 11. TEM substructures for type 317 steel: (a) 1000 °C, 1.0 s-1, dynamic recrystailization grains distinguished by variations in substructure; (b) 1000 °C, 1.0 s-1, dynamic recrystallization nucleus; (c) 1000 °C, 1.0 s-1, dynamic recrystallization grains growing statically into the substructure; (d) 1100 °C, 1.0 s-1, region of static recrystallization grains nucleated around ~ stringers. steel) which induces increased deformation in the surrounding matrix and enhanced nucleation [5, 6, 11, 12]. In addition to having a higher ap and 0 than both type 304 and type 316 steel (see also Fig. 3(b)), it also has

the highest temperature dependence of Op and 0 as shown in Figs. 4-7. The technological significance of the dynamic-recovery-reduced flow stress becomes evident in the calculation of mean pass stresses in rolling which are low

268 I

S-O40"

i

,

i

.7 13

~'E

== so-,~ ~ 10

O,

I-0, C-S,

-.I I

.... IS 16 I / lib 119 P~O sl{ zl2 213 214 EIS El6 LOG(STEADYSTATE FLOW STRESS),LOG 5~ 7b ~o ~o =& ~o EQUIVALENTSTEADYST&TE FLOWSTRESS ~l, MPo

Fig. 12. Plot of the logarithm of the subgrain size against the logarithm of the flow stress for type 317 steel ("), exhibiting a slope o f - 1. os = k d s q

where (~sis the steady state flow stress, k a material constant, d s the subg~ain size a n d q (= --1) the subgrain size exponent. Data for type 304 steel (¢, ref. 19; ~), ref. 20; ~, ref. 21) are provided for comparison.

i

O?

i

i

i

,

,

,

1.4

o+m rE=, 13 1"2 ~oJ T.D. ll .=, .=

I.t z

0.~

~S'O i°,

/md='%-t9 Q ~ - / ÷ ~oJQS ~z

(D

NORMA ,IZAT} N

3"0

5O I@O

~

O-Z 0.1

",-o ,~ ,-,; ~"~

~,o~

h,

1200

,~,

II00

s~,

,;,

I000

,;,'oo ,'+ £% ~, =l= , I~ 16100

,"

Is~

S O 0 +'C

,"

LOGZ (Z" ~I~XP(O/RT),SEC"11

Fig. 13. Plot of the reciprocal subgrain diameter against log Z, exhibiting a linear fit. The subgrains in type 317 steel (.,) are smaller than those in type 304 steel (¢, ref. 19; ~, ref. 21 ; o, ref. 22; @, ref. 23). The two Z scales are shifted to bring the condition 900 °C and 1 s-1 into coincidence [5, 24]. At the point of normalization, d s = 1.11/~m for type 317 steel and d s =1.21 p m for type 304 steel. An activation energy of 400 kJ tool -1 was used for all the type 304 steel data.

in the first pass from the preheat furnace and then increase rapidly in successive stages of the schedule because of the reduction in dynamic recovery as the workpiece cools and the rolling speed increases (the latter being very marked in continuous mills) [5, 24].

The mechanism which would result in the Os* curves in Fig. 2 has been fully explored in aluminium, where it is seen that a polygonized substructure of characteristicdimension and density persistsover strains as great as 60 as a result of a balance between dislocation generation and hardening on the one hand and recovery rearrangement and annihilation on the other hand [11, 14, 29]. The superimposed dynamic recrystallizationcauses further softening through the nucleation and growth of dislocation-freegrains. However, their growth is curtailed,as seen in the formation of repeated boundary necklaces of new grains of final size [11, 13, 30], by the re-establishment within them of the characteristicsubstructure. Moreover, since dynamic recrystallization nuclei are distinguishable only up to two or three subgrain diameters [31], it appears that the cellularityis established in m u c h less strain than in the deformation from the completely annealed state. This probably arises because the small soft nuclei are subjected to intense constraints from their hardened neighbouts. Finally, proof of this contention liesin Figs. 12and 13 where the data points for type 304 steel from both unrecrystallized before the peak (only dynamic recovery) and steady state (both dynamic recrystaUization and dynamic recovery) follow the same relationships relativeto Z and as [19] (Figs. 12 and 13). While in Fig. 6 the narrowing of the gap between ay and a~ as Z decreases is evidence of increased dynamic recovery, a comparison of the curves for type 317 and type 304 steel indicates that the former is m u c h more resistant to softening [6]. In type 317 steelthe athermal region ends at 950 °C whereas in type 304 steel it does so at about 800 °C. At 1000 °C the gap is 90 M P a for type 317 steel and 60 M P a for type 304 steel.However, on heating to 1200 °C the gaps are reduced to 44 M P a and 36 M P a respectively,indicating that dynamic recovery in type 317 approaches that in type 304 steel.This behaviour is comm o n in alloys as precipitates dissolve or solute interactions decrease [24]. Dynamic recovery can also be inferred from the T and ~ dependences of ap (Figs. 1, 4 and 5). The following conclusions also apply to a~* which has been shown to follow the same relationships as ap in type 301 steel [32]. At low Z the power law holds with n

269

suits from the increased significance of glide [33, 34]. Others advocate that the domains of specific mechanisms should not be distinguished but rather that dynamic recovery is associated with a spectrum o f mechanisms which extend across a very broad temperature range for which the exponential law or newer formulations give a more suitable representation [25, 33-36]. The sinh relationship spans the ranges of both the other functions without giving any additional insight into the mechanisms. However, it does facilitate determination of a single activation energy Q which applies across the range of interest [5, 24, 34]. For pure metals at high T, Q is conveniently similar to that for self-diffusion, which confirms the significance of dislocation climb. However, with increased alloying, the value of Q increases markedly from about 280 kJ mo1-1 for pure ~-Fe or mild steel to about 400 kJ mo1-1 for high solute type 304 stainless steel and to about 500 kJ mo1-1 for type 317 steel where precipitation influences the strength at lower temperatures [1-6]. Even higher values of Q are observed in tool steels where the solute effects are less but the change in volume fraction of precipitate is much greater across the range o f normal rolling practice [24]. This is quite a different effect from the variation in activation energy in a pure metal, which usually decreases as temperature decreases and which can be interpreted as the coming into precedence of mechanisms o f lower T dependence [33, 34] or the variation in behaviour of the single mechanism [25, 34-36]. In a simplistic way it can be considered that the increase in Q with increasing alloy content reflects the retardation of the recovery mechanisms by solute and precipitates [5, 11, 24]. From Figs. 1 and 5, it appears that the present results agree with the strengths previously reported for worked type 317 steels [ 7 - 1 0 ] . Two of these steels [9, 10] contain 15-20% more molyb-

5; however, at higher Z the value of n rapidly increases (Fig. 1). At low T or high ~ the exponential law gives a satisfactory fit to the data [5, 6]. The traditional interpretation is that with a power of 4 to 5 the rate-controlling mechanism is dislocation climb; a higher power (or lower strain rate sensitivity m) re-

i

I

I

I

i

i

[

~

i

i

I

4o 3o

2Q

,=

40

120 '160 ' 2 0 ~ 2 4 0 280 $ 2 0 360 4 0 0 4 4 0

80

EQUIVALENT

40 i

80 ,

FLOW

IrrRESl

I~p, tlPiI

120 160 200 240 280 320 360 400 4 4 0 I

I

I

i

i

i

i

0.2 .a 0 3 Z

il

1

0.1

0.4

~- 0.6 ~ OJ

o.a 0.9

w 0 1.0

I.I 1.2

li i ,

,

.

,

xd-h'~,

I'.

I

J ,

,

,I

Fig. 14. (a) Correlation o f ~-o plots for several steels at d i f f e r e n t t e m p e r a t u r e s : e, at intervals o f strain e = 0.1; ], at the critical strain for d y n a m i c recrystallizat i o n ; / , at the end o f a linear segment where subgrain formation commences.

(b) Correlation o f a - e p l o t s for several steels at different t e m p e r a t u r e s : - - - , saturation stress as; +, peak f l o w stress ~p and peak strain ~p. Curve

A s cast or w o r k e d

Steel

1 2 3 4 5 6

As cast Worked Worked As cast Worked Worked

Type Type Type Type Type Type

Reference

317 304 317 304 304 304

[25] [17 ] [ 18 ] [17]

T

(°C) ~ (s-l)

900 704 900 900 900 900

5.0 0.004 5.0 5.0 0.4 5.0

~0 (MPa) 4586 3440 4295 3625 3400 3395

270

denum than the present steel which probably accounts for the higher flow stresses even though one of them has a reduced total solute content and the other has a considerably reduced chromium content. There is a hypothesis that in stainless steels the strength is strongly affected b y the chromium content and very little by the nickel content [1, 2, 5] ; the limited data here do n o t seem to support this. It is also possible that the strength could be increased by a decrease in grain size which is u n k n o w n in the two steels in question. The activation energy has technological significance since, through the Z parameter, it permits calculation of the peak stress or of the mean pass stress for any deformation condition [24]. However, other mathematical formulations are being employed which facilitate more accurate calculation b y c o m p u t e r [ 14] b u t generally lack forms and constants which metallurgists familiarly relate to mechanisms and microstructures. Furthermore, since the strengths of different alloys are fairly similar at a very high T, a high Q indicates a large increase in strength on cooling to the b o t t o m of the hot, working range and hence a difficult alloy to work. The O-a plots show distinctly the effects of dynamic recovery (Fig. 3). Above the limits o f the present graph in stage II, 0 has a high value independent of temperature and strain rate [25, 36]. In stage III the strain-hardening rate decreases at a rate that increases with increasing T and ~ as a result of augmented dynamic recovery [25, 36]. At intermediate stresses the 0-o curves characteristically exhibit a steep linear portion and at high stresses (but fairly low strains) bend to a second linear branch with a lower slope. This usually extends to 0 = 0 in a metal such as aluminium, b u t in type 317 steel, as in many other low stacking fault metals such as copper and nickel [37], there is a final rapid decrease to op as a result of dynamic recrystallization. The shapes of the curves appear to vary in a consistent manner in so far as the divisions between different sections are defined by the three lines slanting upwards from the origin. The lower one designates the critical stress oc for dynamic recrystallization, where the experimental curves diverge downwards from the second linear segment. The other t w o lines demarcate the limits of the linear sections. From Fig. 3, 0 is higher for the cast al-

loy than for the worked alloy because of the presence in the former o f many 8 phase particles which give rise to regions o f higher dislocation density [4, 5]. O-a analysis has been applied to types 301 and 304 steels which are very similar [18, 32]; for any given condition, the 0-a curve for t y p e 304 steel is displaced to the left with a more rapidly decreasing 0 than the curve for t y p e 317 steel (Fig. 3(b)). The change in slope of the individual curves is in some ways remarkable in that the rate o f decrease in 0 is suddenly slowed so that additional hardening and strain take place compared with what w o u l d arise if the curve continued straight (Fig. 14). This change, which appears fairly late in the 0-o curve, occurs fairly early in the a - e curve and has an important effect on its evolution. It might be expected that the change in slope signalled some change in mechanism or microstructure. A study by Carfi et al. [18], from which a single 0-a curve is shown in Fig. 3, discovered that, throughout the straight low a or e segment (high 0), the substructure consists solely of dislocation tangles which increase in density as o and e increase. At e and a near the change in slope, subgralns begin to form, first near the grain boundaries and thence spreading towards the centres o f the grains. With continuing strain, the subgrains become smaller and sharper, b u t the misorientations do n o t change appreciably [18]. Dynamic recrystallization resulted from the formation of new grains near the boundaries, possibly because of bulging. In the very low o or e b u t high 0 domain, extrapolations of the upper linear segments converge to a point on the vertical axis at o0, which was at 4586 MPa and 4293 MPa for the as-cast and worked materials respectively. These values are higher than for any of the other stainless steels tested. For both t y p e 301 and t y p e 304 as-worked steel, 00 was 3395 MPa. It thus appears that 00 is increased by the presence of m o l y b d e n u m which leads to carbide precipitates and in the cast steel to additional 8 ferrite. The dependence of 0s* results from the equation of Kocks and Seeger [34, 36]

'~' = exp{

1-' ln(o'.o*/o's*)} "

eo

RT

J

= l Os* ~ FIsT

\a~o*]

(8)

271 where e0 and F are empirical constants. The activation energy F ln(aso*/a,*) for stage III is considered to be related to cross-slip and is strongly stress dependent [25, 3 4 - 3 6 ] . The saturation stress at 0 K is independent of strain rate. The value of 28 200 MPa is considerably higher than for other stainless steels, e.g. 7500 MPa for t y p e 304 steel [12]. The fact that the behaviour of t y p e 317 steel can be described by a stage III t h e o r y is clear evidence of the significance of dynamic recovery mechanisms during h o t working. The pattern of strain rate sensitivity in Fig. 8 is similar to that in Fig. 1. This was related to power law breakdown, which is consistent with the stage III behaviour above. The convergence of the constant-temperature lines in the log a - l o g ~ plot indicates that, at some high ~, stress is independent of temperature. This totally athermal situation is symmetric with the ~-independent condition at 0 K. The critical strain rate o f about 104 s-1 can be associated with the elastic wave velocity for steel of 5.1 X 105 c m s-1, which represents the upper limit of dislocation velocity. In a further analysis of the results from the viewpoint of low energy dislocation structures as described by Hansen and Kuhlmarm-Wilsdoff [38] in this volume, the relationship of the subgrain size ds to the torsional shear stress r (= 0/3 z/2) was checked against the following equation: ds -

KGb (9)

where K is a structure-related constant, G (= 38.8 GPa at 1100 °C) is the shear modulus and b (= 2.54 X 10 -~ #m) is the Burgers vector. The value of K can be derived from the data points at 1100 °C (0.75Tin) in Fig. 12. From a graphical analysis, K was found to be 12.34 which is in good agreement with that proposed for a low energy dislocation structure. The value is also similar to the K value of 18 determined by Michel et al. [39] for t y p e 316 steel d e f o r m e d in tension between 650 and 816 °C; their lower temperature probably accounts for the higher value of K. In the present calculation the value of K would be approximately halved if the Taylor factor M (= 3.1) were used to convert from o to r. The substructures in Figs. 10 and 11 can also be examined in relationship to the pro-

posal of the low energy dislocation structure analysis t h a t a linear series of subgrains should be misoriented from each other by alternating rotations. The micrographs do show some cases of alternating contrast in agreement with the proposal. The examination of m a n y micrographs shows that the fields of view c o m m o n l y contain a good fraction of cells in dark contrast, although the foil tilt had been selected to make the field uniform in bright contrast with the subboundaries distinctly visible. However, while present, the light and dark contrast is not precisely alternating. In the TEM examination the misorientations between subgrains were frequently checked and found to be about 1 ° and not to accumulate with distance, but alternating misorientation was neither looked for nor noted. Since this is a dynamically recrystallized specimen, although examined after failure at e = 1.25, the strain in the new grains is only of the order of 0.5 (ep in Fig. 2). In conclusion, the substructures observed in the hot working of type 317 stainless steel exhibit the properties of a low energy dislocation structure.

5. CONCLUSIONS In the elevated temperature deformation o f t y p e 317 stainless steel, dynamic recovery gives rise to a polygonized substructure of size defined by Z and related inversely to the flow stress. It also plays an important role in reducing the strength and strain-hardening rate and conferring on both a high sensitivity to temperature and strain rate. While the forms of these dependences are similar to those of t y p e 304, t y p e 317 steel has values of coefficients which make it m u c h stronger t h a n t y p e 304 steel throughout the hot-working range.

ACKNOWLEDGMENTS The authors acknowledge with gratitude the financial support received from the Comitato Tecnologico, Conslglio Nazionale delle Ricerche, Italy, and from the Natural Sciences and Engineering Research Council of Canada. T h e y express their thanks to J. Bowles for assistance in the metallography, to P. McQueen and M. Nagib for drawing the graphs and to

272 I. C r a w f o r d f o r w o r d p r o c e s s i n g . T h e a u t h o r s w o u l d l i k e t o t h a n k P r o f e s s o r G. L ' E s p e r a n c e f o r e n c o u r a g e m e n t a n d t h e u s e o f T E M facilities at the Ecole Polytechnic.

20 21

REFERENCES

22

1 B. Ahlblom and R. Sandstrom; Int. Metall. Rev., 27 (1982) 1-27. 2 W. J. MeG.Tegart and A. Gittins, in J. B. Baltance (ed.), The Hot Deformation o f Austenite, AIME, New York; 1977, pp. 1-46. 3 H. J. McQueen, R. A. Petkovic, H. Weiss and L. G. Hinton, in J. B. Ballance (ed.), The HotDeformation ofAustenite, AIME, New York, 1977, pp. 113-139. 4 H . J . McQueen and N. D. Ryan, Stainless Steels '84, Institute of Metals, London. 1985. pp. 50-61. 5 N. D. Ryan and H. J. McQueen, New Developments in Stainless Steel Technology, American Society for Metals, Metals Park, OH, 1985, pp. 293-304. 6 N. D. Ryan, H. J. McQueen and J. J. Jonas, Can. MetaIL Q., 22 (1983) 369-378. 7 L. A. Norstrom, Scand. J. Met., 6 (1977) 2 6 9 276. 8 C. G. Radu, D. Moisescu, C. Vaida and I. Ilca, Cercet. Met., Inst. Cercet. Met., Bucharest, 18 (1977) 239-246. 9 J. M. Dhosi, L. Morsing and N. J. Grant, in D. A. Edgecombe (ed.), Proc. 4th Int. Conf. on the Mechanical Working o f Steel, Vol. 44i Gordon and Breach, New York, 1965, pp. 2 6 5 - 2 8 3 . 10 W. Roberts, H. Boden and B. Ahlblom, Met. Sci., 13 (1979) 195-205. 11 H. J. McQueen and J. J. Jonas, J. Appl. Metalwork., 3 (1984) 233-241. 12 H. J. McQueen and J. J, Jonas, J. AppL Metalwork., 3 ( 1 9 8 5 ) 4 1 0 - 4 2 0 . 13 C. M. Sellars, in C. M. Sellars and G. J. Davies (eds.), Hot Working and Forming Processes, Metals Society, London, 1980, pp. 3 - 1 5 . 14 W. Roberts, in H. McQueen, J. P. Banon, J. I. Dickson, J. J. Jonas and M. A. Akben (eds.), Proc. 7th Int. Conf. on the Strength o f Metals and Alloys, Montreal, 1985, Vol. 3, Pergamon, Oxford, 1986, in the press. 15 ASTM Stand. E 112, 1982. 16 C. M. Young and O. D. Sherby, J. Iron Steel Inst., London, 211 (1973) 640-647. 17 N. D. Ryan and H. J. McQueen, in S. Blecic et al. (eds.), Proc. Ink Symp. on Plasticity and Resistance to Deformation, pp. 11-26. 18 G. Carfi, C. Perdrix, D. Bouleaux and C. Donadile, in H. McQueen, J. P. Ba~on, J. I. Dickson, J. J. Jonas and M. A. Akben (eds.), Proc. 7th Int. Conf. on the Strength o f Metals and Alloys Montreal, 1985, Vol. 2, Pergamon, Oxford, 1986, pp. 929-934. 19 L. Fritzmeier, M. J. Luton and H. J. McQueen, in

23 24 25 26

27 28 29 30 31 32

33

34

35

36 37 38 39

P. Haasen, V. Gerold and G. Kostorz (eds.), Proc. 5th Int. Conf. on the Strength of Metals and Alloys Aachen, August 19 79, Vol. 1, Pergamon, Oxford, 1979, pp. 9 5 - 1 0 0 . D. H. Warrington, University of Sheffield, personal communication, 1969. M. Ueki and T. Nakamura, Trans. Jpn. Inst. Met., 17 (1976) 139-148. H. B. McShane and T. Sheppard, J. Mech. Work. Technol., 9 (1984) 147-160. M. E. Kassner, A. K. Miller and O. D. Sherby, MetaU. Trans. A, 13 (1982) 1977-1986. N. D. Ryan and H. J. McQueen, J. Mech. Work. Technol., 12 (1986) 279-296. V. K. Kocks, J. Eng. Mater. Technol., 98 (1976) 76-85. H. J. McQueen and W. B. Hutchinson, in N. Hansen, A. Horsewell, T. Leffers and H. Lilholt (eds.), Deformation o f Polyerystals: Mechahisms and Microstructures, Proc. 2nd Ris@ Int. Symp. on Metallurgy and Materials Science, September 14-18, 1981, Ris} National Laboratory, Ris~, 1981, pp. 335-342. D. J. Abson and J. J. Jonas, Met. Sci., 4 (1970) 24 -28. J. R. Cotner and W. J. McG.Tegart, J. Inst. Met., 97 (1969) 73-79. H. J. McQueen, O. Knustad, N. Ryum and J. K. Solberg, Scr. MetaU., 19 (1985) 73-78. R. Sandstrom and R. Lagneborg, Scr. Metall., 9 (1975) 59-65. H. J. McQueen and H. Bergeson, Met. Sci, 6 (1979) 25-29. N. D. Ryan and H. J. McQueen, in H. McQueen, J. P. Bailon, J. I. Dickson, J. J. Jonas and M. A. Akben (eds.), Proc. 7th Int. Conf. on the Strength of Metals and Alloys, Montreal, 1985, Vol. 2, Pergamon, Oxford, 1986, pp. 935-940. W. D. Nix and B. Ilschner, in P. Haasen, V, Gerold and G. Kostorz (eds.), Proc. 5th Int. Conf. on the Strength o f Metals and Alloys, Aachen, August 1979, Vol. 3, Pergamon, Oxford, 1979, pp. 15031530. H. J. McQueen and H. Mecking, in B. Wilshire and D. R. J. Owen (eds.), Proc. 2nd Int. Conf. on the Creep and Fracture of Engineering Materials and Structures, Swansea, April 1-6, 1984, Vol. 1, Pineridge, Swansea, 1984, pp. 169-184. W. Blum, H. Munch and P. D. Portella, in B. Wilshire and D. R. J. Owen (eds.), Proc. 2nd Int. Conf. on the Creep and Fracture of Engineering Materials and Structures, Swansea, April 1-6, 1984, Vol. 1, Pineridge, Swansea, 1984, pp. 1 3 1 148. H. Mecking, in M. F. Ashby (ed.), Dislocation Modelling o f Physical Systems, Pergamon, Oxford, 1981, pp. 197-211. R. E. Cook, G. Gottstein and V. F. Kocks, iT. Mater. Sci., 18 (1983) 2650-2664. N. Hansen and D. Kuhlmann-Wilsdorf, Mater. Sci. Eng., 81 (1986) D. J. Michel, J. Moteff and A. J. Lovell, Acta. Metall., 24 (1973) 1269-1277.