Effects of alloying upon the hot workability of carbon, microalloyed, tool, and austenitic stainless steels

Effects of alloying upon the hot workability of carbon, microalloyed, tool, and austenitic stainless steels

Journal of Mechanical Working Technology, 12 (1986) 279--296 279 Elsevier SciencePublishers B.V., Amsterdam- Printed in The Netherlands EFFECTS OF ...

965KB Sizes 2 Downloads 10 Views

Journal of Mechanical Working Technology, 12 (1986) 279--296

279

Elsevier SciencePublishers B.V., Amsterdam- Printed in The Netherlands

EFFECTS OF ALLOYING UPON THE HOT WORKABILITY OF CARBON, MICROALLOYED, TOOL, AND AUSTENITIC STAINLESS STEELS

NORMAN D. RYAN and HUGH J. McQUEEN

Department of Mechanical Engineering, Concordia University, Montreal H3G 1M8 (Canada) (Received October 15, 1984; accepted December 7, 1984)

Industrial S l i m ' m a r y Through torsion simulation of industrial hot working, the characteristics during and after deformation of carbon, microalloyed, tool and 7 stainless steels have been related to their austenitic microstructures. In the 900--1100°C, 0.1--10 s -1 range, the peak flow stresses depended on strain rate exponentially and temperature b y an Arrehenius term. Generally, strain-hardening rate, peak stress, activation energy and strain for dynamic recrystallization" increased with rise in either solutes or precipitates. The hot ductility decreased as alloying increased and the temperature and strain rate decreased, except for a maximum at 1.0 s -1 for stainless and tool steels. Dynamic and static recrystallization decreased the propagation of fissures b y isolating them from grain boundaries.

Introduction The hot forming of steel represents a very large tonnage so that an improvement in processing leads to an important reduction in costs. The scientific study of hot workability has progressed significantly in the past two decades; however, the transfer of information to metalworking engineers has not apparently proceeded to the most beneficial level. The purpose of this paper is to summarize and compare the hot-forming characteristics of low carbon, HSLA, tool and austenitic stainless steels in a format which emphasizes industrial relevance to a higher degree than in their initial publication as research reports [1--5]. The properties of similar and related steels drawn from the published literature are also included to confirm the behaviors observed by the authors. The differences in behavior are related not only to compositions but to microstructures in order to develop a more widely applicable understanding.

0378-3804/86/$03.50

© 1986 Elsevier Science Publishers B.V.

280

Mechanisms of deformation and softening The basic mechanisms which provide the reduced flow stress and high ductility in elevated-temperature shaping have been fully identified [6-17]. The strain hardening which accompanies dislocation flow in cold processing is increasingly reduced, as the temperature T is raised or the strain rate ~ is lowered, by dynamic recovery through dislocation annihilation and arrangement into low energy sub-boundaries. In addition, at higher strains (~ 0.7, 50% reduction) in austenite, continual dynamic rec~ystallization eliminates dislocation substructure by the repeated passage of highangle grain boundaries. The strain necessary for nucleation can be accumulated over several stages of forming. Both before and during dynamic recrystallization, a substructure with dimension d s characteristic of the conditions becomes dynamically stable and determines the flow stress, which is inversely proportional to it [6--16]. Solutes in general reduce both the rate of dynamic recovery through inhibiting dislocations from leaving their slip planes to undergo annihilation and the rate of dynamic recrystallization through lowering the mobility of the grain boundaries to which they have segregated. Precipitates also slow down both recovery and recrystallization, the first by hindering dislocation motion and stabilizing dense substructures and the second by pinning grain boundaries. Dynamic recovery improves the ductility by reducing the formation of w-cracks at triple grain-boundary junctions through enhancement of accommodating grain flow by reducing lattice strength relative to grainboundary sliding resistance. On the other hand, dynamic recrystallization slows propagation of intergranular cracks originating at either inclusions or triple junctions by displacing the boundaries from the cracks which causes them to arrest and blunt until they once more become associated with a boundary [3, 7--14, 17]. Through curtailing dynamic recovery and recrystallization, alloying additions raise the lower boundary of the hotworking temperature range. Alloying also lowers the upper limit by decreasing the solidus temperature. Thus as solutes and second-phase particles increase, the working range diminishes until the incidence of failure makes hot-shaping unfeasible [3, 10, 13, 17]. In addition to the dynamic mechanisms there are alsa static restoration mechanisms which take place during intervals between passes or in final cooling [1--5, 8--14, 17]. Static recovery can lower the attained hightemperature flow stress by as much as 20% and if it alone occurs, the substructure is preserved and either affects the flow stress in the following stage or increases the strength of the product. Static recrystallization usually takes place quite rapidly in austenite, the rate increasing with pass strain, strain rate and temperature, and the grain size decreasing with only the first two. The grain size affects either the workability in the subsequent pass, or the properties of the product. The behavior of the present steels in this regard will be the subject of a subsequent paper.

281

Experimentaltechniques The steels (Table 1) on which the authors experimented, were of commercial origin and had been rolled to plate with the exception of the stainless steels which were as-continuously-cast [1--5]. While the properties of as-cast steels are significant in the break-down stages, laboratory specimens can only reflect the microsegregation and not the macro
Low C Medium C 0.032 Nb 0.050 Nb 0.050Nb A2 H13 M2 M2 T1 304 304 304 304 316 316 317

0.14 0.42 0.09 0.12 0.12 1.00 0.38 0.84 0.86 0.72 0.07 0.05 0.07 0.06 0.017 0.07 0.035

C

0.56 0.94 1.54 0.94 0.94 0.70 0.30 0.25 0.31 -1.76 0.92 1.76 1.15 1.84 1.27 1.73

Mn

0.009 0.25 0.22 0.007 0.007 0.30 1.00 0.30 0.30 -0.68 0.11 0.82 0.53 0.52 0.64 0.44

Si

0.06 0.88 . . 0.04 0.04 5.00 5.25 4.00 4.05 4.13 18.38 18.2 17.6 18.16 16.92 17.2 18.6

Cr

0.09 0.16 . 0.04 0.04 ---0.20 -8.68 11.3 8.52 9.36 12.42 10.9 13.88

Ni

0.012 0.02 . 0.007 0.007 1.15 1.30 5.00 4.85 0.40 0.08 0.02 0.27 0.22 2.76 2.92 3.22

Mo

0.002 0.002 0.30 1.05 1.90 1.78 1.04 --(0.17") --(0.32*) (*Co)

---

V

--( 6 . 5 0 ÷) ( 5 . 9 8 ÷) ( 1 8 . 1 6 ÷) ---(0.01") --(+W)

--0.032 0.050 0.050

Nb

C o m p o s i t i o n s o f c a r b o n , H S L A , t o o l a n d s t a i n l e s s s t e e l s ( w e i g h t %)

TABLE 1

0.005 -----

---0.30 --

------860 -300 110 ---

45 45

--

48 80 66

N (ppm)

--

--

--

--

Ti

6.45 7.60 17.25 17.92 23.73 27.07 30.57 29.16 29.58 32.10 32.77 37.87

--

--

0.93 2.26 --

S a n k a r et al. Barraclough, Sellars Ouchi and Okita S a n k a r et al. K n u d s e n et al. I m b e r t et al. I m b e r t et al. I m b e r t et al. Carlsson, Roberts Coward R y a n et al. Barraclough, Sellars M c Q u e e n et al. Mase R y a n et al. H u g h e s et al. R y a n et al.

Metallic R e f e r e n c e solute

5 5 5 23 30 4 24 26 27 4 22 4

2

1

1 24 25

No.

G¢ b0

283

""' ' 420

140

360

120

Ib ~; 3 0 0

I00

NO' STIrEL ~:l's" T~'C ' REFERENCe" • o M2 b 31(5 • C.317 HI3 ,d A2 f *304 • g'316

o Q. ~ ~ d ~

• k M2

.-= 80

~

~g

" ~

Z

i

etaL et~al.al/''~"

SortkarRyanKnUdsenet



h

Imbert ,t a,.

:]/ o.* ~=

' L.

60

" |

1.0:

*3=7 : .o.3O4

~

60 (~

Imbert et al. Ryan et ol.

900

zIO 'm HI3A2

i 180

5.0

Imbert et el. Hugho$ et al. Ryan et el.

....... rl•0 Jl*304h 0.14C0"05 Nb:

.~

~ 240

5,0" 7.0

.oo

,

Ryo. ,, 0,

* AS'CAST MATERIAL

..

I

20 0

I 5 •

I I I I0 15 20 NUMBER OF TURNS |

l

i

l

!

I 25

I 30

i

2

4 6 8 I0 12 14 EQUIVALENT' STRAIN, Fig. 1. Representative high temperature flow curves for C [1], HSLA [1, 2], stainless [4] and tool steels [5]. The strength of 304, 316 and 317 increase in order of solute concentration, of A2, H13 and M2 in order of increasing carbide content. The curve for 316 was added for comparison [22].

and the strain rate declines (from 5.0 to 0.1 s -1, Fig. 1), the flow stress decreases and the ductility increases. The highest strengths are those of the tool steels which at 900 ° have 39 vol% carbide in M2 and 10 vol% in H13 and A2. They are much softer at 1100°C because the carbides are then only 15, 1 and 0.3 vo]% respectively (but then the solutes are 8.4, 7.6 and 6.5 respectively) [5]. The stainless steels are next in strength with 317, 316 and 304 having solute levels of 38, 34 and 29% [4, 22]. The solute has the double effect of decreasing both the stacking fault energy and hence the dynamic recovery and also the mobility of the grain boundaries and hence the dynamic recrystallization [4, 6--11, 14, 15]. There are other influences however; the 316 has the lowest C content and the 317 has 3.2% Mo, which has partly precipitated. At 1100°C, 0.1 s -1, the stainless steels are stronger than the tool steels in which the carbides have dissolved. The as-cast stainless steels exhibit low ductility because they contain considerable 5 ferrite [3, 4, 7, 13]. The strength data from a wider selection of steels [1--5, 22--24] are presented more clearly in Fig. 2. The peak stresses are a reasonable basis of comparison since, although peak strains vary a little, they give an idea of how high the stress is in a pass of 50--60%. It can now be seen that

284 I

I

I

I

STEEL "~-M2 AJ • 316 • ~17 Hl3-1r~ U A2 . . _ J • 304 • "M)4-~

..... ..... 400 .....



316

0

a.

316_1

00.05Nb E] 0.14 C

I

I

REFERENCE Corlsson ,Roberts Imbert et ol. Hughes el at. Ryan et ol. Imber1' (11 al. Bar roctough,Seltars

Ryan et at. Knudsen et al. Sankm ~ at.

%

Ib

- 300

%

GO

M2 %

o_1 -0

~ 2oo

"%%

~

W a. z

o w

I00

STRAIN RATE, ~ II.Os-~ (ExCept where noled~ I

I

i

I

'

i

900

950

I000

1050

I100

I1,50

TEMPERATURE,°C

Fig. 2. The equivalent peak stress decreases as the t e m p e r a t u r e increases (~ = 1 s-l). The o r d e r o f increasing strength is C, H S L A , stainless and t o o l steels [ 1 - - 5 ] . Between 900 and 1000°(:; almost all the carbide dissolves in A2 and H13, leaving t h e m softer than the stainless. The strength of 316 at 3.9 s -1 is s o m e w h a t higher t h a n the present results [ 22 ].

dependence on structure and composition is greatest at the lowest temperature and least at the highest, either because precipitates dissolve, or solutes have less influence on dislocation glide or recovery. To say the same thing in another way, the stronger materials b e c o m e progressively stronger as temperature declines. More than other steels, H13 and A2 strengthen markedly b e t w e e n 1000 and 900°(:; because the carbide content jumps from about 1 to 10 vol% [5]. The 317 is considerably stronger than the 316 and 304 because of greater solute content and carbide formation [4, 1 3 ] . The alloy steels are clearly much more resistant to deformation than the low-carbon steel; a relatively miniscule micro-alloy addition has considerable effect on the strength. In the as-cast stainless steels, 0.5--5 #m particles of 5 ferrite cause strengthening by locally increasing the strain hardening of the matrix [4, 7].

285 The peak strains (Fig. 3), which for practical purposes can be considered the critical strains for dynamic rec~stallization, are higher for the stronger steels by 50--100% than for C steel (but not in all cases) [1--5, 22, 24, 25]. The rise in peak strain' does make a contribution to the augmentation in flow stress. However, in other cases, it has been argued that the increase in strain-hardening rate advances nucleation as is seen for 304 and 317 stainless steels [4]. It is also seen that microalloying raises ep markedly, indicating that it is through this effect that it raises the strength [1, 2]. Solute~,,and precipitates retard recrystallization through exerting a drag on migrating boundaries in the same manner as they retard static recrystallization [7--16]. I

I.

I

I

I

I

0.9 S T R A I N RATE, ~ = 1.0 s -I (EXCEPT WHERE NOTED)

0.8

,,~ 0.7 t

"

.-_,

0 0.4 Z ~n* 0 . 3 I-0")

0.~

0

-..

-

M2

/

~A-~H t3 - 304

-

~ ~ STEEL CONDITION REFERENCE,,~,,,,m ~ 0 0.05Nb rolled L ~ Knudsen el el. " ~ " 317 -0 003Nb rolled Ouchi and Okito ---- • 304 rolled Borreclouoh end Sellore - ' - ~1 extruded Hugh.. e t e h ,,rj 316 us-cost Ryen et el. --&)

OI

o.o,N,



--I-I I

900

M2] A2 A HI3

.

roiled 1 forgedJ

3041

Imbert et o l

317 J

as-cast

Ryon et el.

0.14C

rolled I

I

Sonker el al, I

I

I

950

I000

1050

I100

1150

TEMPERATURE, °C Fig. 3. The strain t o the peak, which is closely related to the critical strain for dynamic reerystallization, decreases as T increases, ep generally increases as alloy content rises; however, the order differs from that for a s particularly for 317 and 304 where high rates o f strain hardening give rise to early nucleation [1--5]. Data on 316 [ 2 2 ] , 304 [24 ], 0.03Nh [25 ] give support to the present results.

286

Stress, strain-rate, temperature dependence Figure 4 illustrates a commonly-used relationship between stress o and strain rate $ with empirical constants A and ~ [1--7, 23]. =

A(T)exp(flo)

(1)

The slopes fi of the curves are characteristics of the materials and independent of the temperature. When the temperature is increased, the intercepts A decrease so that the curves shift w i t h o u t change in slope. The value of fi is seen to decrease from 6.1 to 3.9 as the strength of the steel rises, which indicates an increasing strain-rate sensitivity. The cause for this is essentially the same as that of the higher temperature dependence appearing in Fig. 2. The value of the equation is t h a t stresses for any strain rate can be calculated once the empirical constants are determined. |

I

I

I

I

I

1

"~= A(T) exp (,~)

I050Oc

~.o

Z

/~

0.5

//,/?/7 i/'

~ 0.~F

/

i

| l

--

I /

I / /

I

I

'

" ,;

/./ Jf

,'///

.,/ /

/,

I-/'/J /.7'[

/

/I[I OI;////•

I

/

I

I I/

/ / /

....

I// •1/ #

I I

i I00

i 140

.

005 Nb Knudlen Ill 016 4 504 McQueenet al.&2 .

.

41' 304 5 Ryon el ol. 5.4 " •

IO00 ~,. except where noted i 60

. 0 e

. . . . , ~_,,,berte~ol. 5.e A HI3J 5.7

.....

/

I

/

/

STEELRE~E~'E~E~_~.~ a,4c s~°wo,.,o,. ~.,

i leO

51?.J (*as cost)

....



M2

~ i 220



M2 i 260

48

Corisson 5.9 °nd Roberl$ Irnbet"lId OI. 3.9 i 300

EQUIVALENT FLOW S T R E S S ~ ,(MPo)

Fig. 4. F o r a given material, a graph o f log ~ versus o generally gives a series o f parallel straight lines for different t e m p e r a t u r e s ; here o n l y o n e line is s h o w n for each alloy. The hardest material has the smallest value of/3 and hence the highest strain rate sensitivity [ 1 - - 5 ] . F o r comparison, additional results for M2 and 304 are included [23, 2 6 ] .

287 '1 ' 36C 16

~t4 ,~w

42C

---•30C

360

~ -~

--

• -

~ '

I

' I ' TEMPERATURE STEEL REFERENCE RANGE V M2 Corlsuon "~IISO-IOSO*C ~ d RobertsJ ~100-1000 eC • M2 Imbert el M.-~tO00-900*C • 317 * 1 Ryon I t oL I , • 3 0 4 , J(*AS-CAST)JIO00-SO0 C 00.OSNb KmNIIin et at11 ,

[] o.,4c

'"

I '

=3 ooo-,oo

S0.k0,.,

0.03 Nb Ouehl ond

--.~

"6

111

c

I

I' QHW kJ mo 436 435 7IS 502 40Z 434

sos 4

~ ~1 ' STRAIN RATE ~ = 1.0 s"1

300

P,1,6

&t

24C

OI

J ~

/.~ ~

i/

J~"

14 _

240

~



g

.o,o E

18C

Bil I0

" '24C ~.180

o 8

~ .180 ~

~

~/"

' 120

120~1

~

4•

-"*

:t

.......

! .60

~

I60

~ ~

. s, Io ' c

0.70

• .oo'c I 0.72

, o=5. o . c

0.74

0.76

,ooo'c , = 0.78

0.80

95o-c h 0.82

900°C 0.84

8

60

_.

,o /

oi

0.86

I / T x 1 0 3 , °K'=

Fig. 5. Graphs of/~a versus 1/T exhibit a linear relationship for each steel. The hardest material has the highest temperature dependence with slope proportional to QHW. Supporting evidence for M2 and 0.03% Nb steel [23, 25 ]. TABLE 2 Comparative values of n, ~ and activation energy QHw Steel

n

~ QHw Reference (10 -2 MPa -1) (kJ tool -1)

No.

Fe--0.14C Fe--0.12C--O.05Nb Fe---0.09C--0.03Nb A2 0--1Mo--5Cr--3V--1C H13 0--1Mo--5Cr--IV--0.4C M2 6W--5 Mo--4CY--2V--0.8C H13 0--1Mo--5Cr--1V--0.4C M2 6W--5Mo--4Cr--2V--0.8C 304--18 C2---9Ni--0.08Ms 316--17Cr--12Ni~2.76Ms 317--19Cr--14Ni--3.22Mo 304--18Cr--8Ni--0.27Mo 304--19Cr--9Ni--0.22Ms 304--18Cr--11Ni--0.02Mo 316--17 Cr--11Ni--2.92Mo

5.1 8.6 5.3 9.2 9.5 9.7 8.0 -4.3 4.5 4.0 4.3 4.4 5.8 4.7

6.1 6.4 6.4 5.9 4.9 3.9 4.9 3.9 5.4 5.4 4.8 5.2 5.3 4.7 5.6

1 2 25

303 435 401 395 401 435 424 438 407 401 508 410 393 410 460

Sankar et al. Knudsen et al. Ouchi and Okita

Imbert et al.

5

Samanta 36 Carlsson and Roberts 23 Ryan et al. as-cast material

4

M c Q u e e n et al. Ouchi and Okita Barraclough, Sellars Hughes et al.

T h e t e m p e r a t u r e T d e p e n d e n c e is s h o w n i n Fig. 5 [ 1 - - 1 4 , 2 3 , 2 5 ] w h i c h i n c o n j u n c t i o n w i t h e q n . (1) p r o v i d e s t h e e x p r e s s i o n = A'exp(~a)

exp(----QHw/RT )

(2)

26 25 24 22

288

with A' and QHW empirical constants and R = 8.3166 J/mol °K. The intercept A'exp(flo) evidently increases with stress, the curves being parallel. The slopes of the straight lines are apparent activation energies QHW {Table 2) and indicate a thermally-activated atomic process. The curves are generally straight over this narrow temperature range indicating a single rate-controlling mechanism; however, it has not been possible to specify it since the value is usually higher than that for self-diffusion or for dislocation creep [7--10, 14]. Although microstructural studies confirm dynamic recrystallization, QHW is generally larger by 10--50% than the value for static recrystallization [4--7, 11, 14]. The activation energy increases with alloy content, reflecting the rising difference in flow stress between 1100 and 900°C; one can conjecture that QHW rises because the alloying makes the operation of the restoration mechanisms, i.e. dislocation climb, crossslip or node unpinning and grain boundary migration, more difficult. The change in slope for the tool steels reflects the rapid increase in hot strength between 1000 and 900°(3 as carbides precipitate [5]. For most of the steels, i

STEEL

' i

220I ~

240

, P

-~

~

3,7

| ~f p ~ M ; )

~E 180 V. # nO

(tension)

tl /.*

tb"

"~,~

~.--~."~.... 7

M2

°C

E,s-'

900

,,,0 ~,

4.;)x,o 2.

0.4,

I000

5xlO 0

55x1021

040

,,0o

5,Io°

LT,IO" 0.SZ I.;) xlO 17

031

,ooo

,,,o-,

,.o,,o,°

o.44

,,,oo 5,,o0

2.s.,o,6 9o,,o,

0.50 o.4,

15xlO 16

0.42

"~' " " "" "" 1150 "~_

"....,, M 2

IOxlO0

HI3

140 I100

5xlO 0

950 |U/f"

~ n w ~ . ~

I00 |l

Nb 1050

~ 9 5 0 ~ " ~ . . . _ ~

80

60 (:~ w

y 40

_

20 O

0

0.5

ISxlO 16

0.6; )

3.4,10'~ 0.50

IxlO 0

6.4 x I017 0.49

0.42 C

950 1050

9.9x10 -2 2.7x1012 I.OxlO 0 2.7x1012

0.55 0.52

0.42 c

950 1050

I l x l O -2 2.7x10 N 99x10 -2 26x10 II

044 0.48

Ryon et ol. Imber! l ! 0|. Carillon and Roberts Borroclough and Sellors Knudsen et ol. i i

1.0 1.5 EQUIVALENT

o.,,

IxlO-'

,,lo-'

REFERENCES

' 1,2,8 5.5,6.7 4 9,10,15,14,15,16 11,12 i

,.,,,o,,

I000

15 I"zLLI .J

Z,s -i

7 = AeJ~e -Q/RT ~-e+Q/RTZ = -Ae':" ~*o = AI %0 I/2 Z n I

2.0 STRAIN,

Fig. 6. F l o w curves for several steels selected so that the value o f Z is t h e same for each alloy. The m a t c h is bad for M2 and p o o r for H 1 3 p r o b a b l y because o f different carbide dispersions [ 1 - - 5 ]. S u p p o r t i n g data for M2 and 0 . 4 2 C [ 2 3 , 24 ].

289 calculating flow stresses at temperatures above or below the experin~ental range is justified as long as there are no phase changes. Thus extrapolation to lower temperatures is not possible for the tool steels, nor for the C or HSLA steels when they become ferritic. It is worthy to note that the presence of ~ ferrite in as-cast stainless steel does not seem to influence the activation energy. The writing of eqns. (1) and (2) with ~ on the left is essentially a format for creep, where they were first used [6, 7 ] . They can, however, be rearranged into a format which combines the independent variables into a single parameter Z (Zener--Holloman) [ 1 - - 1 1 ] : o = A" + log Z = A" "+ log [~ exp(+Q/RT]

(3)

This equation states that a is constant when ~ and T are both altered to maintain Z constant; this is illustrated in Fig. 6 where the stress--strain curves for different temperatures are the same because the Z condition is the same [ 2 4 ] . As a result of eqn. (4), a material's hot strengths for a variety

0.14 C

14oo 13oo i ,O,Ol

ip,

0.05 Nb

no M2

,2oo

i,oo

1200

'lO0

'P" l,

'P" I

,ooo

I ,p!'

1 'P"

11°''

I000

'P"

I'P"

110,,

1020~ L/N

1021 M2

j 3o7 "

HI3

~

..~

A2

304

/.,,,,'~fj

.I "="" " ="

/ /1 ~ ' ~ So.ko.., oL.,,,,,, ,~""~..,,,,~'~,,,Jw~

0 AZ

--

240

800

lip '`

900 I0 tg I

Z (~,exp~,O/RT) NORMALIZED AT 900"C TEMPERATURE FOR DEFORMATON AT lOs "1 ,1~ 420 STE,EL REFERENCE o3 • M2 Corlsson ond Robert1 , • M2 Imbert et al. 360 • ~17 ] / ~ I :304 | *Ryon et OI, ~ • zl6 J Io, co,t1 ,,/ ..... 304 Oorrocloughand Seltors ~ 300

~u. ~

900

o o,.c

~

(1.

0.05 Nk

180

~

120

~

6O

-r (.9

-r 317 304 316

w'-

ho"

0.1 ,,-'

,~"1

16'°

1300

I~" I

,~'~

LOG

15"

1000

1200

'4)'~Io,3o,6"1.'oo

I

1,6"

I100

' t o "IZOO [

HI3, A2

/

I

16"1

11oo

Z,

s -I and

I

I

,6"

i ,o"l

900

I~"

I100

900*C I0 s-I

1.0 s-I t

lib"

Io;ol

I000

,@' ,ooo I T°C

gO0

,6"

1,6" ,o0

IOl~) 80 vj

16"

sooli

for I.Os "t

Fig. 7. When the peak f l o w stresses o f a material are plotted against log Z the data f r o m different temperatures lie on the same line [1--5 ]. The lines for stronger alloys have an increasing slope (due to higher QHw). The Z scale has been s u p p l e m e n t e d by temperatures for ~ = 1 s -I . In addition, the horizontal scale has been normalized to 900°C, 1.0 s -1 chosen because it is near the lower working limit. Data for M2 [ 2 3 ] and :}04 [ 2 4 ] confirm the present results.

290

of temperatures and strain rates lie on a single curve when plotted against Z (Fig. 7) [1--5, 23]. This makes interpolation easier once Z is calculated for the forming conditions desired. Since calculating Z is a little clumsy and requires looking up QHW, a temperature scale for ~ = 1 s -~ has been attached to each Z scale; this is different for each alloy since the value of QHW varies. For strain rates which are multiples of 10, a temperature value can be located by displacing the appropriate number of decades on the Z scale. Because the value of QHW differs with the alloy, when stresses are plotted against a c o m m o n log Z scale, the curves are displaced along the horizontal axis in a rather disconcerting manner. To preserve the manifestation of alloying effects from Fig. 2, it was decided to normalize the curves at a c o m m o n condition: 900 °, 1 s -~, which was selected because important differences in strength occur there. This places the curves in the correct vertical order; however, it still leaves a mismatch in the temperature scales since t h e y are stretched more for higher values of QHW" In Fig. 8, log ep is presented as a function of Z and results in a single curve for each alloy [1--5, 11, 24]. Just as the order of ep in Fig. 3 did n o t

Fe 0 3 4 % C

1400

Fe 0 . 0 5 Nb M2

0

I00

090 -Or

080

-

07o

z

~ -02

~-

' " 0 60

z

1300

I~'00

1200

I100

I0 Is

I0 le

STEEL

n

• 316 * 0 O05Nb

0081 0059

• u~

0063 1

Z~ HI3

0064

0 A2

I100

REFERENCE Knudlen

et el

900

I000 I018

I017

*Ryon et o l

I000

~

.

oo5o

• 3,7



00.3

900 1 0

1IoZO9

1021

Z([exp(+glRT)) NORMALIZED AT 900°C TEMPERATURE FOR DEFORMATION i I -AT l O s /

,..,4"316

I Imberf et al,

I

0066J

,30,

800

/,Aw/

I /.," R, . . . . .

,

Do.c [~ 0 4 2 C

o~96 0L89

so.w., ,, o,. 8arraclough

--.sL.

0,.3

s.,o,.

O05Nb

/

( . . . . . . . ,)

~ ~

/ j

L.:./~--~

~

~ , , , ~

(~,,'~'~l~'~- ~"

...-'rp:~..~..q~

"

-- 304

....-"~-

w

_o, -06

030 o25

316 H I 3 , A~'

o.2~//

tO

1300

I;)O0

/o,.c

I100

LOG Z, s"~and T°C

900oc OIs I

I000

i

900oc I0 s '

900

900oc I0s-

BOO

for I.Os-'

Fig. 8. For each alloy, log ep versus log Z is a straight line; the order of alloys differs from Fig. 7. The lines for C and HSLA steels have unusually high slope. Supporting data for HSLA steels [11] and 0.42C steel [24].

291 agree with the order of Op in Fig. 2, the order here is different from Fig. 7. In general, the curves do not cross, indicating that the difficulty of nucleating dynamic recrystallization relative to other alloys remains the same across the range studied. The outstanding exception to this is the C steel where the different slope in Fig. 3 is emphasized. As evidence that this is not a discrepancy due to one set of results, data for a 0.42% C steel give a parallel line [24]. Data for a compilation of HSLA steels [11] are included and show a slope half way between the C steel and the other alloys. The relationship for the HSLA steels was determined to be [ 11] : ep = A'" dg0 ~zn

(4)

The plot of Fig. 8 confirms eqn. 4 since initial grain size (dg0) differences would shift the curves along the horizontal axis without changing the slope. The grain-size dependence was not examined in the present case since grain sizes are known only for the tool steels. The importance of dg0 depends on the role of the grain boundaries as the sole site for nucleation of dynamic recrystallization in most cases. It is not known how the presence of carbides or 5 -ferrite particles affect the nucleation.

Hot ductility The dependence of high-temperature ductility on temperature and strain rate are shown in Fig. 9 [1--5, 22, 26--33]. In Fig. 9(a), the ductilities are generally for strain rates greater than 1 s -1, but for the C and HSLA steels they are given for 0.1 s -1 since at 1 s -1 they would be above the present graph; this emphasizes the greatly reduced ductility of the alloy steels compared to C steels. The ductilities of almost all the steels increase with temperature over the range examined. These increases are the result of the acceleration of dynamic recrystallization with rise in temperature, which increases its effectiveness in retarding the propagation of grainboundary fissures [3, 7, 9, 14, 17, 20, 26, 31, 32]. Studies over a wider range show that, as temperature declines, the ductility diminishes to a minimum where grain-boundary fissuration is unimpeded due to the absence of dynamic recrystallization. In austenitic stainless steels, ductility rises at still lower temperatures as fracture becomes transgranular when grain-boundary sliding ceases to occur [7, 13, 20, 26, 31, 32]. In steels that transform on cooling, the ductility declines further as the e grains form at the 7 grain boundaries, thus impeding their migration and leaving intergranular cracking unchecked; however, as the volume of ferrite exceeds 50%, the ductility rises until the transformation is completed, and then declines down to ambient temperature [7, 9]. On the other hand when the temperature is increased, ductility increases to a maximum and then declines as the solidus temperature is approached due to liquation at the grain boundaries [3, 7, 9, 17, 32]. Apparently, the decline in ductility

292

of the A2 steel is a result of liquation since it has the highest C content and could also be affected by segregation and the presence of trace impurities; the H13 and M2 also showed this behavior at other strain rates (Fig. 9(b)) [5]. The tool steels have the highest C contents and lowest melting temperatures o f the alloys discussed. The as-cast stainless steels have very poor ductility compared to wrought stainless steels [5, 2 7 ] . This arises because the ~ ferrite particles seriously hinder the migration of the 7 grain boundaries even though they are softer [3, 4, 7, 9, 24, 27]. The ductility usually rises with strain rate to a maximum and then declines [3--5, 7, 9, 17, 31--33]. Increasing the strain rate from about 10 -3 s -1 usually accelerates dynamic recrystallization because it creates a denser dislocation substructure and hence a greater driving force for recrystalli!

r

! CONDITION

STEEL -

-

m m

16

....

r-I 0.14 C -~ 0 0 0 5 NbJ H 13 t 321 • 304

roiled

Sankar et al. Knudeen~et oh

forged

Irnber~et

~

I.o 12 n," IJ-

o I- I0

(a

ol.

rolled

0 A2

-14

i

Nich61son et oh rolled Mc~ueen et al. J rolled f m b e r t et oh

& M2

t.=.l n'* :E)

i REFERENCE

--~-- • __._,_ • - •

316 extruded 304 s l o b "i 316 ] ingot J

----



304|

- -



3~7 j

as-cost

Hughes et el. Mase Ryon et ol.

z o~ 8

1-O3 t"Z .J

/

6

~=5 s-i ::)

o w4

= 3.9 s-' ~=4.2 s-;

2

~C m

~=ss"

f:=4.zs"= 0.13 $-I

0

I

900

I

1

I

950 I000 1050 TEMPERATURE, °C

I

I100

293

(b]

18 DNF IO00"C

/

16

14

~

30

'

STEEL

U

u_ 2S

t-U

e.,

(/)

" (0

zi,v

R

|

i ii ~

A2

, , , , , , , ~ , A HI3

& az

~ ~ McQueen et ot =• 5 1 6 , ~ = • 3 t 7 ~ V Ry~m st ol,

~'~

• 3 0 4 - - I - G i t l l n s ot ol. (tension)

o ,~,, '} 15

Imbert et Or.

J

.....

I.-20

z

REFERENCE

0 0.14 C 5mkor et el. 0 0.05 I ~ Knudsen et el.

900%

~s i0

4

%

j*

'--"--"B------,----,,-,-

,10-z s 0 •

@

,

IO00"C

G056

5. (~1 STRAIN

I.O

-I0.0

RATF.. ~ ~ t "l

Fig. 9. The ductility of the steels increases with rise in T (a) with the only exception being A2 which diminishes at the highest T because of liquation near the solidus [1--5 ]. The ductility usually increases with ~ because dynamic recrystallization is speeded up (b). A peak and decline occur where the increased flow stress o f the grains enhances grain b o u n d a ~ cracking and propagation. Complimentary data on 304 [26, 27], 321 [28, 29], 316 [22], T1 [30] and 304 in tension [33].

zation. However, the critical strain for dynamic recrystaUization also rises (as does e~, Fig. 3), so that there is more opportunity for w-cracks to initiate and for mcreased cracking at inclusions [~, 7--9, 17, 19, 26, 32--35]. The height of the maximum and the strain rate for it depends upon the material. In Fig. 9(h), the results for the C and HSLA steels lie on the ascending branch of the curve and the maximum o c c u r s at a high ~ because these alloys have almost no solute or precipitates [1--&]. For the H13 tool steel which, in general, exhibits very high ductility, the maximum occurs about 1 s -1 at both 900 and 1000°C [5]. For the ot~er-tool steels and the

294

as-cast stainless steels, the data lie on the descending branch of the curve since the maxima were at 0.1 s -1 or lower [4, 5]. Multistage processing: estimation of stresses In industrial hot-shaping operations, it is common to have a series of passes which are usually at successively lower temperatures and higher strain rates [7--12, 14, 17]. The flow curves of several steels for a hypothetical schedule consisting of 5 equal passes of 18% reduction (e i = 0 . 2 ) , during which T decreases from 1100 to 900°C and ~ increases from 0.1 to 10 s -1, are presented in Fig. 10(a) [1--5]. The curves illustrate h o w solute and precipitate raise the processing forces for alloys, most markedly at low temperatures and high strain rates. Since these are derived from isothermal tests on specimens which for each material were preheated in the same way before cooling back to the deformation temperature, they are only valid if the schedule permits complete recrystallization to the original grain size between each pass. In addition, calculation cannot show that cracks generated in the earlier passes do not propagate in the later ones to 450

STRESS STRAIN CURVES FOR ROLLING PASSES DETERMINED OR INTERPOLATED FROM ISOTHERMAL TORSION DATA TORSION- INTERRUPTION STRAIN, ; = 0.2 SIMULATED ROLLING-REDUCTION, R ~ 18%

-

400

g. ~ 350

,b"

~ 300 hi

STEEL

REFERENCE

M2 * 517 ]

Imbert

* 316 J

Ryon et al.

A2

Imbert el ol.

et al.

25O

f

/

0.050 Nb Knudsen el ol. 0.14 C Sankor el at. *AS-CAST MATERIAL

p-

/

S ,¢I----:

M2 .317

200 ~Z

,y

.316 A2

,oo

o,.:

150

0.05 Nb

%

IlO0°C, 0.1 s -I

%

1 0 5 0 ° C , 0.3 s -I

STAND 4

J STAND 3 L

950°C,

I O 0 0 ° C , 1.0 s -I

EQUIVALENT

STRAIN,

i%.

STND

5

3 . 0 s-I 9 0 0 = C , I0.0

E:

Fig. 10. F l o w curves for a multistage schedule with declining T and rising ~ compiled from isothermal continuous-flow curves; static recrystallization during intervals [1--5 I. Tool and stainless steels become progressively stronger than the C steel as T decreases.

295 cause failure before the schedule is complete. T he problems o f carry over o f grain and subgrain structures and o f fissures f r o m one pass t o the n e x t can most easily be resolved by simulation o f t he schedules; such simulation is discussed in a subsequent paper. Conclusions Th e h o t workabilities o f C, HSLA, stainless and t ool steels have been described and c o m p a r e d over t he c o m m o n range for h o t forming. The flow stress, th e strain-rate sensitivity, the t e m p e r a t u r e d e p e n d e n c e and t h e susceptibility t o fracture were shown t o increase with rise in solute and precipitate levels in a m a nne r qualitatively similar t o their effects on cold strengthening and annealing resistance. The use o f Z t o extrapolate or interpolate values of flow stress was confirmed and a new plotting m e t h o d i n t r o d u c e d t o simplify t he p r o c e d u r e and make comparisons between alloys easier.

Acknowledgements The authors wish to express their gratitude t o the Natural Sciences and Engineering Research Council o f Canada and t o the FCAC program o f Quebec for th e financial assistance which made this series o f researches possible. The collaboration o f Atlas Steels and o f SIDBEC--DOSCO is also acknowledged. Finally, recognition must be given for the helpful discussions with o u r colleague Prof. J.J. Jonas o f McGill University. References 1 J. Sankar, D. Hawkins and H.J. McQueen, Met. Tech., 6 (1979) 325--332. 2 W. Knudsen, J. Sankar, H.J. McQueen, J.J. Jonas and D. Hawkins, in C.M. Sellars and G.J. Davies (Eds.~, Hot~Wb~rl~iitg and~Formirll~Processes, ThG Metals 'Society, London, 1979, pp. 51--58. 3 H.J. McQueen, J. Sankar and S. Fulop, in K.J. Miller and R.F. Smith (Eds.), Mechanical Behavior of Materials (ICM3), Pergamon Press, Oxford, 1979, Vol. 2, pp. 675--684. 4 N.D. Ryan, H.J. McQueen and J.J. Jonas, Can. Metall. Q., 22 (1983) 369-378. 5 C. Imbert, N.D. Ryan and H.J. McQueen, Metall. Trans., 15A (1984) 1855--1864. 6 J.J. Jonas, C.M. Sellars and W.J. McG. Tegart, Met. Rev., 14 (1969) 1--24. 7 C.M. Sellars and W.J. McG. Tegart, Int. Met. Rev., 17 (1972) 1--24. 8 C. Rossard, The Microstructure and Design of Alloys (ICSMA 3), Inst. of Metals, London, 1973, Vol. 2, pp. 175--203. 9 H.J. McQueen and J.J. Jonas, in R.C. Arsenault (Ed.), Treatise on Materials Science and Technology, Plastic Deformation of Materials, Vol. 6, Academic Press, New York, 1975, pp. 393--493. 10 W.J. McG. Tegart and A. Gittins, in J.B. Ballance (Ed.), The Hot Deformation of Austenite, AIME, New York, 1977, pp. 1--46. 11 C.M. Sellars, in C.M. Sellars and G.J. Davies (Eds.), Hot Working and Forming Processes, The Metals Society, London, 1979, pp. 3--15.

296 12 H.J. McQueen, Can. Met. Q., 21 (1982) 4 4 5 - 4 6 0 . 13 B. Ahlblom and R. SandstrSm, Int. Met. Rev., 27 (1982) 1--27. 14 W. Roberts, in G. Krausz (Ed.), Deformation, Processing and Microstructure, ASM, Metals Park, Ohio, 1983, pp. 109--184. 15 T. Sakai and J.J. Jonas, Acta Met., 32 (1984} 189--209. 16 H.J. McQueen and J.J. Jonas, J. Appl. Metalworking, 3 (1984) 233--241. 17 H.J. McQueen and J.J. Jonas, J. Appl. Metalworking, 3 (1984) 410--420. 18 H.J. McQueen and J.J. Jonas, in A.L. Hoffmanner (Ed.), Metal Forming, Interrelation Between Theory and Practice, Plenum Press, New York, 1971, pp. 393-428. 19 H.J. McQueen, in G. Sih and C.J. Provan (Eds.), Defects, Fracture and Fatigue, Martinus Nijhoff Pub., The Hague, 1983, pp. 459--471. 20 S. Fulop, K. Cadien, M.J. Luton and H.J. McQueen, J. Test. Eval., 5 (1977} 709-714. 21 G.R. Canova, S. Shrivastava, J.J. Jonas and C. G'sell, in J.R. Newby and B.A. Niemeier (Eds.), Formability of Metallic Materials-2000AD (STP 753}, ASTM, Philadelphia, 1981, pp. 189--210. 22 K.E. Hughes, K.D. Nair and C.M. Sellars, Met. Tech., 1 (1974} 161--169. 23 J. Carlsson and W. Roberts, End cracking in connection with hot working of highspeed steel (Research Project 4357/79), Jernkontorets Forskning Rep. D 401, Nat. Inst. for Metals, Stockholm, 1982. 24 D.R. Barraclough and C.M. Sellars, Met. Sci., 13 (1979) 257--267. 25 C. Ouchi and T. Okita, Trans. Iron Steel Inst., Japan, 1982, Vol. 22, pp. 543-551. 26 H.J. McQueen, R. Petkovic, H. Weiss and L.G. Hinton, in J.B. Ballance (Ed.), The Hot Deformation of Austenite, AIME, New York, 1977, pp. 113--139. 27 T. Mase, Sumitomo Search, 3 (1970) 31--38. 28 A. Nicholson, Iron and Steel, 37 (1964) 290--294 and 363--367. 29 A. Nicholson, D. Smith and P. Shaw, Deformation Under Hot Working Conditions (Special Report No. 108), Iron and Steel Institute, London, 1968, pp. 161--166. 30 M.D. Coward, Ph.D. Thesis, University of Sheffield, Sheffield, England, 1970. 31 A. Gittins and W.J. McG. Tegart, Metals Forum, 4 (1981} 57--62. 32 A. Gittins, J. Aust. Inst. Metals, 20 (1975} 184--200. 33 A. Gittins, L.G. Hinton and W.J. McG. Tegart, Manufacturing Eng. Trans., 2 (1973) 199--207. 34 J.A. Charles, in C.M. Sellars and G.J. Davies (Eds.), Hot Working and Forming Processes, Metals Society, London, 1980, pp. 87--98. 35 S.L. Semiatin and J.J. Jonas, Formability & Workability of Metals, Plastic Instability & Flow Localization, ASM, Metals Park, Ohio, 1984. 36 S. Samanta, Deformation Under Hot Working Conditions (S.R. 108), Iron and Steel Institute, London, 1968, pp. 122--130.