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The effect of nickel on the beginning of transformation of austenite in a 0.55 carbon, 0.35 molydenum… Scott, Donald Alexander 1947

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t i l l fry  THE EFFECT OF NICKEL ON THE BEGINNING OF TRANSFORMATION OF AUSTENITE IN A 0.55 CARBON, 0.35 MOLYBDENUM STEEL  by DONALD ALEXANDER SCOTT  A THESIS SUBMITTED IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE DEGREE OF  MASTER OF APPLIED SCIENCE IN THE DEPARTMENT OF  -  MINING AND METALLURGY  * * * * * *  THE EFFECT OF NICKEL ON THE BEGINNING OF TRANSFORMATION OF AUSTENITE IN A 0.55 CARBON, 0.35 MOLYBDENUM STEEL ABSTRACT This study i s carried out to f i n d the e f f e c t of n i c k e l on the beginning of Isothermal transformation of austenite i n an iron-carbon-molybdenum a l l o y containing 0.55 carbon and 0.35 molybdenum.  An introduction describes the  iron-carbon equilibrium system, the products of slow c o o l i n g of austenite, the r e l a t i o n between slow cooling and isothermal transformation at temperatures below equilibrium, and a f u l l discussion of terminology used.  A l i t e r a t u r e review discus-  ses theories of transformation and previous work on the e f f e c t of n i c k e l on austenite transformation.  Development of experi-  mental technique i n Isothermal transformation and melting of pure alloys i s discussed. The Isothermal transformation diagrams are shown f o r beginning of transformation of austenlte of base composition 0.55 carbon and 0.35 molybdenum, r e l a t i v e l y f r e e from impurities ( s i l i c o n , manganese, e t c . ) , showing the e f f e c t of n i c k e l on the beginning of transformation. used are 0, 2*13, 3.69, and 5.31 percent.  Nickel additions Isothermal trans-  formation i s shown by photomicrographs which are discussed fully. The e f f e c t of increasing n i c k e l on the isothermal  t r a n s f o r m a t i o n o f an a l l o y c o n t a i n i n g 0.55 c a r b o n and 0.35 molybdenum I s as f o l l o w s : delayed appreciably:  (1)  the p e a r l i t e r e a c t i o n i s  (2) f e r r i t e formed a t I n t e r m e d i a t e  temperatures (880 t o 1000 d e g . P . ) becomes more  prominently  a c i c u l a r , the a c i c u l a r f e r r i t e r e a c t i o n t a k i n g the place o f t h e upper b a i n i t e r e a c t i o n o f l o w n i c k e l a l l o y s :  (3)  the  a c i c u l a r f e r r i t e r e a c t i o n i s f o l l o w e d f i r s t by r e j e c t i o n o f c a r b i d e p a r t i c l e s , and l a t e r by a g g l o m e r a t i o n the c a r b i d e phase:  (4)  and growth o f  t h e a c i c u l a r f e r r i t e and f e a t h e r y  b a i n i t e r e a c t i o n s as r e p r e s e n t e d  on t h e i s o t h e r m a l t r a n s f o r -  m a t i o n diagram become s e p a r a t e d by t h e appearance o f a bay i n the i s o t h e r m a l t r a n s f o r m a t i o n  curve.  T A B L E  OF  C O N T E N T S Page No.  OBJECT  1  SUMMARY  2  INTRODUCTION  4  THEORIES OF TRANSFORMATION  11  THE EFFECT OF NICKEL  16  EXPERIMENTAL TECHNIQUE IN MELTING INDUCTION FURNACE  23 23  GLO-BAR FURNACE  .24  EXPERIMENTAL TECHNIQUE IN ISOTHERMAL TRANSFORMATION  29  EXPERIMENTAL RESULTS  38  DISCUSSION OF RESULTS  62  CONCLUSIONS  67  ACKNOWLEDGMENTS  69 * * * * * * * * *  1. THE EFFECT OF NICKEL ON THE BEGINNING OF TRANSFORMATION OF AUSTENITE IN A 0.55  CARBON, 0,55 MOLYBDENUM STEEL  OBJECT The phenomena accompanying the decomposition of austenite have not yet been  f u l l y or s a t i s f a c t o r i l y explained  despite the fact that they have been examined by many observers and that they play an extremely important part i n the heat treatment of s t e e l .  E s p e c i a l l y lacking i s the e f f e c t of  single elements added to a l l o y s of i r o n and carbon i n which no other residual elements are present to complicate the experimental r e s u l t s .  Once this e f f e c t of i n d i v i d u a l elements  has been established It should be possible to f i n d the e f f e c t of one a l l o y on that of another either d i r e c t l y or i n d i r e c t l y , and, by so doing, some of the phenomena accompanying austeni t e decomposition may be explained more s a t i s f a c t o r i l y than they are at present.  Studies of this type should provide  quantitative data f o r r e l a t i n g the Isothermal transformation curve and hardenability, and from which the form of t h i s curre and the hardenability curve could be predicted. E s t a b l i s h i n g data f u l l y f o r the above involves a long range study, and t h i s thesis describes only one step toward the ultimate objective. The object of t h i s phase of the work Is to establ i s h the quantitative e f f e c t of n i c k e l on the beginning of  2. Isothermal transformation of austenlte i n a pure iron-carbonmolybdenum a l l o y containing 0.55 carbon, 0.35 molybdenum, and substantially no s i l i c o n , manganese, or other a l l o y i n g element.  SUMMARY A method of melting has been established f o r producing pure a l l o y samples having a predetermined  carbon and  a l l o y content, substantially free from s i l i c o n and manganese. Base materials are Norway i r o n , pure graphite powder, pure f e r r o - a l l o y s , and pure n i c k e l shot. The method of melting involves using approximately 100 grams of material i n a 35 cc alundum c r u c i b l e , protected against thermal shock and oxidation In a globar furnace by a shaped graphite block container and l i d .  The melts are  poured as quarter-inch-diameter Ingots i n a graphite mold. The ingots are annealed 30 minutes at 1500 deg.P. and cut into one-twelfth inch discs before Isothermal heat  treatment.  The beginning of isothermal transformation of austenite has been determined f o r a l l o y s containing 0.55 carbon, 0.35 molybdenum, and 0, 2.13, 3.69, and 5.31 percent n i c k e l at temperatures  650, 700, 750, 800, 880, 930, 1000,  1100, and 1200 deg. P. f o r times up to 20,000 seconds (5^ hours).  A photomicrographic record was made of beginning of  transformation, and of progress of transformation at a l l temperatures where two products form. Adding increasing amounts of n i c k e l to an a l l o y containing 0.55 carbon and 0.35 molybdenum gives the following effects on the isothermal transformation of austenlte: (1)  the Isothermal transformation curve i s pushed generally  to the r i g h t , (2) the p e a r l i t e reaction i s delayed much more than any other reaction, (3) the tendency towards formation of a c i c u l a r f e r r i t e at Intermediate temperatures i s more noticeable, and on growth of a c i c u l a r f e r r i t e a carbide precipitates from f e r r i t e , agglomerates, and grows as a separate phase formed d i r e c t l y from austenite, (4) the upper bainite reaction i s displaced by the a c i c u l a r f e r r i t e reac0  t i o n , (5) the lower b a i n i t e reaction as represented on the curve becomes separated from the rest of the curve by the appearance of a bay between the a c i c u l a r f e r r i t e and lower bainite reactions, and (6) the upper temperature of formation of lower b a i n i t e i s not lowered as rapidly as the eutectoid temperature.  4. INTRODUCTION To understand thoroughly p r e c i p i t a t i o n of the various products from austenite during cooling or at constant temperature  I t i s necessary to have a clear understanding of  the basic metallurgy of pure a l l o y s , or the i r o n - r i c h portion of the system iron-carbon.  The portion of the iron-carbon  constitution or equilibrium diagrami r e l a t i n g to p l a i n carbon alloys and t h e i r heat treatment i s shown i n Figure 1. For any s t e e l of the carbon range covered i n Figure 1, heating f o r several minutes to the area marked austenite at a temperature above the l i n e s A 3 and Acra transforms the s t e e l e n t i r e l y to austenite, a s o l i d solution of carbon i n face-centered (gamma) i r o n .  By slowly cooling any s t e e l of  carbon content less than 0.8 percent (eutectoid or p e a r l i t e composition) a s o l i d solution of carbon i n body-centered (alpha) iron of extremely low carbon content, c a l l e d f e r r i t e , begins to precipitate at approximately the indicated by the l i n e A 3 .  temperature  Further slow cooling to the temper-  ature marked A^ p r e c i p i t a t e s a l l the excess or proeutectoid f e r r i t e , and, as the f e r r i t e p r e c i p i t a t e s , the composition of the remaining austenite progresses with decreasing temperature down l i n e A 3 to the eutectoid composition (0.8 percent carbon). The proportion of f e r r i t e to austenite i s shown i n Figure 1 for a 0.4 carbon s t e e l at 1400 deg. F. as the r a t i o of Y 2 to 1 - Subscripts r e f e r to the bibliography at the end.  1800-  PSficeNT  Fi'f I  - Iron-Car6e» Steel  Heat  CM  BON  £cju///6r/vm Di'qg ram Treatment  Portion.  Showing  6. x y. As the temperature i s slowly lowered below the A i l i n e the remaining austenite transforms to eutectoid ( p e a r l i t e ) , which has a microstrueture  of alternate p a r a l l e l cementite  (Fe$G) plates and f e r r i t e .  No appreciable change i n struc-  ture occurs between A^ and room temperature. Cooling any s t e e l above eutectoid composition i s similar to the above except that the proeutectoid  constituent  appearing at the l i n e Acm i s cementite, a compound of iron and carbon with preferred orientation of I n t e r s t i t i a l s o l i d s o l u t i o n carbon atoms i n a body-centered i r o n l a t t i c e . The transformation  process f o r heating a s t e e l i s  the reverse to that described f o r cooling. Figure 1 i s reproduced i n practice only with steels of low r e s i d u a l a l l o y content, and then only with extremely slow cooling, which may be expected from the f a c t that the l i n e s on the diagram represent equilibrium between two phases which requires that the free energy i s the same f o r each phase.  P r e c i p i t a t i o n of any phase from austenite can occur  only when the free energy of the phase Is lower than that o f the austenite, the excess free energy of the austenite that transforms being used up i n surface eneryy of the p r e c i p i t a ting phase.  The adjustment i n composition of the remaining  austenite changes the equilibrium conditions to give the austenite a free energy equal to that of the p r e c i p i t a t i n g phase.  P r e c i p i t a t i o n stops when the free energies are equal.  p.—  100%  or Acm  MARTENSITE LOG  TIME  Fig2- Schematic Representation cf Isothermal Trans  8. Cooling at any rate faster than that required f o r oonditions given i n f i g . 1 changes equilibrium r e l a t i o n s h i p s and produces a difference In the microgtructure of the transformed product, and i f cooling i s f a s t enough metastable phases other than f e r r i t e and p e a r l i t e are produced.  These  d i f f e r e n t conditions are best related In f i g . 2 which i s a schematic representation of the isothermal transformation diagram.  Temperature i s represented v e r t i c a l l y , time on a  logarithmic h o r i z o n t a l s c a l e .  The two s o l i d curved l i n e s  represent beginning and end of isothermal transformation of austenite a f t e r quenching r a p i d l y from a u s t e n i t i z i n g temperatu<re to isothermal temperature.  The dotted l i n e represents a  proeutectoid constituent, f e r r i t e f o r hypoeutectoid  steels,  cementite f o r hypereutectoid s t e e l s , which i s non-existent for a eutectoid s t e e l .  The area between the proeutectoid l i n e and  the beginning curve i s greater the farther the carbon content of the s t e e l i s from the eutectoid composition.  The two h o r i -  zontal l i n e s Ms and Mf represent 0 percent and 100 percent martensite which forms during cooling i f the cooling i s f a s t enough to avoid the curved portion, and also during cooling i f transformation i s arrested at any time after beginning at a temperature higher than Ms by quenching the specimen to room temperature; the martensite reaction i s not isothermal. Martensite has a highly strained, d i s t o r t e d  body-centered  l a t t i c e structure which approaches a tetragonal l a t t i c e .  The  iOQVcM A= fiustenite h^THcomfotiti^n  f .  Fi$3"  5%M 2S%3  7  S0%*1 S0%5  7S%&  Joo%6  Product  T<Me  at Temperature T,  Method of  Studying  MotatboraphtcaUy  Izothtrmat  Transformation  foprese/rftd  ScAe/nat'Co/ty  10. body-centered for  l a t t i c e i s attained a f t e r reheating martensite  tempering. A schematic method of producing the isothermal  transformation diagram i s shown i n f i g . 3. The products of transformation of s t e e l vary with temperature  of formation.  Just below the A-^ temperature  p e a r l i t e which forms i s coarse lamellar, almost spheroidal. As the temperature  of formation i s lowered the r e s u l t i n g  p e a r l i t e lamellae are f i n e r and the s t a r t o f the reacti on i s increasingly f a s t e r down to -the "nose" which represents the maximum rate of reaction.  At this maximum another very  f i n e l y lamellar product i s formed, commonly c a l l e d upper bainite.  Below the "nose" the product i s b a i n i t e , an acicu-  l a r aggregate of extremely f i n e feathery lamellae, now more commonly referred to as lower b a i n i t e .  The austenite which  does not transform to martensite on quenching and holding isothermally at a temperature  between Ms and Mf eventually  transforms to lower b a i n i t e . P a r t i a l l y transforming to any product and quenching to room temperature  immediately stops the reaction and as the  specimen i s cooled through the range Ms to Mf the untransformed austenite transforms t o martensite.  This has made  possible isothermal studies of the beginning and progress of transformation by metallographic methods.  II. THEORIES OF TRANSFORMATION During, the past few years several theories have been advanced to explain the mechanism of nucleation and growth of p r e c i p i t a t i n g phases during austenlte decomposition In s t e e l s .  The f i r s t work reported on isothermal transfor-  mation i s that of Bain and Davenport2, and from their work has been developed what i s now known as the isothermal transformation diagram. Later investigators^ ^have shown that the bottom portion of "S-curve" should be a series of horizontal l i n e s , i n d i c a t i n g that martensite does not form isothermally, but . only by lowering the temperature. have r e a l i z e d the importance  Recent i n v e s t i g a t o r s ^ g 5  of i n t e r p r e t i n g the curved  portion of the '^-curve" as two "C-ourves , namely, p e a r l i t e 11  and bainite, although previous workers recognized a d i f f e r ence i n the type of transformation to p e a r l i t e and b a i n i t e . A t h e o r e t i c a l basis f o r the Isothermal transformation diagram was presented by Zener  &  which explains some of the l e s s e r  understood f a c t s , but t h i s has been open to question  7 Q  ^  because of too many assumptions which may not be v a l i d . Hultgreng recently proposed s i x d i s t i n c t reactions i n h i s description of decomposition  of austenlte i n a l l o y s t e e l s .  An experimental f a c t that transformation to pearlite and bainite normally begins at grain boundaries has been shown to be t h e o r e t i c a l l y true by Zener^*  With t h i s assumed and  12. observed f a c t i n mind e a r l i e r w o r k e r s ^ ^ ^ - ^ p r e s e n t e d a n a l y t i c a l theories of nucleation and growth of p e a r l i t e , a l l based on certain assumptions.  I t must be noted that any  theory f o r s o l i d - s o l i d reactions cannot be accepted as f a c t because many of the variable factors cannot be measured or, rather, have not yet been measured.  Yet many facts can be  observed e x p i r i c a l l y even i f they cannot be d i r e c t l y measured, and a n a l y t i c a l theories may  be produced by analogy  with those experimentally measurable for g a s - l i q u l d - s o l l d reacti ons • Mehl and J e t t e r . ^ , Mehl and Bowman have advanced 1Q  the nucleation theory that, during the s o l u t e i  14  the random migration of  atoms i n austenite, volumes of the correct com-  p o s i t i o n and s i z e are established by chance, depending upon energy r e l a t i o n s h i p s , for the formation  of a nucleus,  and  this t h i s nucleus can remain or grow only I f the free energy o f the p r e c i p i t a t i n g phase i s lower than that of austenite. Wells and M e h l , P e l l l s s i e r , Hawkes, Johnson and lg  Mehl  16  have shown that l o c a l d i f f u s i o n currents of solute  atoms are very necessary f o r the formation and growth of p e a r l i t e , keeping i n mind that f o r p l a i n carbon s t e e l s only carbon and iron atoms have to  d i f f u s e , whereas i n a l l o y  steels each a l l o y atom, as w e l l as carbon and i r o n , has to diffuse either to or away from a p o t e n t i a l nucleation s i t e . Mehl  i n  and Ham, advanced the theory that d i f f u s i o n away from 7  13. an area of a p o t e n t i a l carbide nucleus of an a l l o y element not appearing i n the carbide phase and d i f f u s i o n towards that area of elements appearing i n the carbide phase must take place simultaneously before the p o t e n t i a l nucleus can form. Mehl  1 Q  related h i s d i f f u s i o n nucleation theory to  interlammellar spacing decreases to describe the gradual change from p e a r l i t e to b a i n i t e nucleation at the p e a r l i t e maximum or "nose" of the  M  S-curve . n  The greatest number of l a t t i c e d i s c o n t i n u i t i e s or imperfections exist where there i s a break i n curvature of the l a t t i c e which i s the case f o r grain boundaries,  the  greater the curvature or smaller the radius the greater the discontinuity of the l a t t i c e .  For impure steels inclusions  also cause sharp d i s c o n t i n u i t i e s .  Zener4 postulated that  a l l o y i n g elements tended to congregate at l a t t i c e tions (although why  imperfec-  they should i s questionable) and the  carbide or f e r r i t e which forms usually has d e f i n i t e high or 1 ow a l l o y content compared with normal p l a i n carbon s t e e l f e r r i t e and carbide. Because a l l o y i n g elements change the carbon sion rate very  l i t t l e  i4 i6,i8,19  a n d  a  l  l  >  o  y  diffu-  elements d i f f u s e  at much lower rates than does carbon^ , nucleation and growth 7  i s delayed by the slower d i f f u s i o n of the a l l o y i n g element. N  This same type of reasoning can be applied to the delay i n formation of proeutectoid phases arid b a i n i t e .  14. Some general and s p e c i f i c a l l o y effects upon the "S-curve" of austenite decomposition Davenport2Q  a n <  have been presented by  * B a i n ^ , but i t i s assumed because i t i s not g  stated otherwise and because manganese i s included i n the analysis of steels used, that a l l o y additions have been made to commercial s t e e l s .  Much work has been coordinated by the  United States Steel Corporation and published i n t h e i r "Atlas of Isothermal Transformation"  gg  giving the isothermal trans-  formation diagrams of many commercial s t e e l s .  An extensive  compilation of "S-curves" published up to 1945 has been presented by M o r r a l , nearly a l l of which are f o r impure or g 3  commercial a l l o y s . Some workers have noted..  1 Q  0  .  the p a r t i t i o n of  O A  a l l o y i n g elements between the f e r r i t e and carbide phases, and have advanced some theories on the p a r t i t i o n e f f e c t to explain the e f f e c t of a l l o y i n g elements upon the "s-curve".  The main  outline of t h i s theory i s that some alloys appear only In f e r r i t e , others only i n carbide, and some with varying degrees of p a r t i t i o n between the two phases.  The low d i f f u s i o n rate  of a l l o y elements slows down formation of n u c l e i of s u f f i c i e n t size to grow, and also slows down growth by the slow movement of the various p a r t i t i o n i n g elements to the respective phases. Investigators have i n d i c a t e d ^  1 Q  2  5,26 27  b  y  x  *" y ra  studies and other means the theory that p e a r l i t e i s nucleated by cementite  or carbide, and b a i n i t e i s nucleated by f e r r i t e .  isBainite has the same orientation r e l a t i o n as f e r r i t e i n respect to the austenite from which I t i s formed.  Bainite  d i f f e r s from proeutectoid f e r r i t e only i n microstrueture and composition, the difference In composition being that b a i n i t e has a higher Carbon content on formation and retains increasingly more carbon to completion of transformation formation  temperature i s lowered.  as the  One possible further proof  that proeutectoid f e r r i t e and b a i n i t e are s i m i l a r l y formed from the same nucleus i s offered by Hollomon, J a f f e , and Norton 3» i n showing that holding a hypoeutectoid 2  isothermally i n the proeutectoid range without  steel  transforming  v i s i b l y to f e r r i t e , markedly reduces the time to form b a i n i t e by a subsequent quench to the b a i n i t e temperature.  This i s  even more revealing when i t i s shown that the f r a c t i o n a l times (to form v i s i b l e amounts) held i n each range add up to unity. This would substantiate the theory that nucleation begins the instant the s t e e l i s reduced below the c r i t i c a l temperature »  range.  The apparent additiveness o f the f e r r i t e and b a i n i t e  reactions i s possibly not general since energy relationships and d i f f u s i o n rates are d i f f e r e n t at d i f f e r e n t temperatures, and the nucleation and growth rates would depend on energy and d i f f u s i o n r a t e s . The theory that nucleation begins i n s t a n t l y on an atomic scale i s contrary to the theory that nucleation s t a r t s on a microscopic  scale about v i s u a l beginning, as i s usually i  j  16.  assumed i n theoretical d i s c u s s i o n s 4 ^ o , l l , l 2 , 2 9 *  T b  -  e  s  a  m  e  variance of opinion was voiced i n t r y i n g to explain p r e c i p i - , tation hardening where hardening Is obtained before p a r t i c l e s could be detected by X-ray, which leads to an assumption  that  "quasi-nuclei" can possibly e x i s t . Some of the above theories can be combined or expanded to give the theories presented by Blanchard, Parke and HerzlggQ, and Bowman^ that decomposition of austenite can be explained on a basis of rates of nucleation and growth of both cementite and f e r r i t e , each of the four rates being some function of temperature. The reaction to martensite i s a d i f f e r e n t type, i n which a new set of energy relationships supplant the normal nucleation and growth reactions, and some form of s t r a i n mechanising 3 4 10 ° P  e r a t e s  occurs as the temperature  3  0  that instantaneous reaction  i s lowered, and the s t r a i n thus set  up by p a r t i a l transformation changes the required energy d i t i o n s so that a lowering of temperature  con-  Is necessary to  cause further transformation.  THE EFFECT OP NICKEL Some previous investigators,„ „_ „ „„ have * 17,19,21,22,32,33 n  noted the e f f e c t of n i c k e l on the Isothermal transformation of various types of s t e e l s , most of which were of commercial grade with varying amounts of impurities such as manganese  17. and s i l i c o n .  This did not permit a s t r i c t l y comparative,  quantitative estimate of the e f f e c t of n i c k e l .  I t has been  reported by DavenportgQ that, i n general, n i c k e l moves the w  S-curve  w  to the r i g h t with l i t t l e change i n shape except that  caused by lowering the Ae temperatures. Both n i c k e l and manganese produce a sluggish transformation reaction i n carbon,free  iron a l l o y s and lower con-  siderably the Ae Temperatures i n carbon a l l o y s sluggishness was  3 1  34.  reported.^ to be so great that two  This years  were required to measure equilibrium i n a 4 . 8 percent alloy.  nickel  As a r e s u l t of these observations Hami7 assumes that  carbon does not necessarily have to be present i n manganese and n i c k e l alloys of Iron to r e t a r d the decomposition of the gamma phase (austenite i n carbon a l l o y s ) .  Hence Ham  concludes  that manganese and nickel retard austenite transformation because of their presence In the o r i g i n a l austenite and i n the subsequently-formed f e r r i t e which must contain both elements. This i s apart from any e f f e c t they may have on the r a t e of formation and growth of any p r e c i p i t a t i n g phase other than ferrite.  Expanding the reasoning further, the e f f e c t of  manganese and n i c k e l i n s h i f t i n g the p e a r l i t e reaction to longer times may be caused by the fact that they must appear in f e r r i t e rather than i n the carbide phase, and that t h e i r inherent sluggishness on the austenite to f e r r i t e reaction delays the austenlte to carbide reaction proportionally.  o  z  r~i$ 4 - Effect  +  *  e  to  it.  of Nickel on Eutectoid Carbon Content inSM  19. Ham^y has recognized that d i f f u s i o n must be accounted f o r i n any theory explaining the decomposition of austeni t e by nucleation and growth.  Much useful data has been  presented on d i f f u s i o n c o e f f i c i e n t s f o r n i c k e l ^  3 g  ,  manganese , c a r b o n ^ y , and molybdenum-^^ . 35  14  3  38  Ham's explanation of the e f f e c t of n i c k e l and manganese on the "s-curve" does seem to hold true In the b a i n i t e region where i t has been shown i s formed from f e r r i t e n u c l e i .  gg  g  g ^ that b a i n i t e g  Because both n i c k e l and  manganese produce an equilibrium sluggishness In the gamma to alpha transformation the delay i n proeutectoid f e r r i t e and b a i n i t e reactions should be proportional to the degree of sluggishness i n the gamma to alpha  transformation.  This reasoning may not hold true f o r the effect i n . delaying the p e a r l i t e reaction since i t i s generally conceded that p e a r l i t e i s nucleated by c a r b i d e by f e r r i t e .  4 1  Q  ^  ^  and not  The common conception i s that the addition o f  carbide-forming elements (Cr, Mo, W, V, T i ) delays nucleation and growth of carbides simply because o f d i f f u s i o n requirements of the carbide-forming elements, and that n i c k e l additions with carbide forming elements delay the r e a c t i o n s t i l l more because of sluggishness.  However, t h i s may not be  the case as i t i s conceivable that the stronger carbideforming elements may hasten nucleation so that extremely small "quasi-nuclei" form, much smaller than the stable size f o r  Sain Vareopott  V  1700  \  x  \  N  \  \  V \  •  \ \ V  N  N \  1100  \ \  \ \  • -  \  \ \  Ni ckel  900  t  \  4 e 'o 'a f/g 5-Effect of Nickel on Eutectoid Temperature in o  *  *  21 growth requirements, yet p o t e n t i a l l y stable.  The delay i n  transformation may be caused by the slow d i f f u s i o n of the "quasi-nuclei"  to form nuclei of a size required by energy  relationships  f o r growth to occur and large enough to see.  It i s also conceivable that ferrite-formlng  elements con-  tribute to the delay i n the p e a r l i t e reaction because of i n t e n s i f y i n g effects on the carbide-formers. Graphical representation of the effects of a l l o y i n g elements on the Ae temperature and the eutectoid position i s given by Baing^.  carbon com-  Reasonable agreement up to 6  percent n i c k e l has been shown by Marshy f o r the e f f e c t of n i c k e l on the eutectoid  carbon content, which Is reproduced  i n f i g . 4 with the addition n i c k e l on the eutectoid  of Bain's data.  The e f f e c t of  temperature i s shown i n f i g . 5.  Fty 6 - Crucible For Induction Melting-  23.  EXPERIMENTAL TECHNIQUE IN MELTING After preliminary experimental work melting s t e e l samples In a high frequency induction furnace using various types of crucibles and i n a glo-bar type furnace, i t was decided that the most s a t i s f a c t o r y melts can be obtained i n the  glo-bar furnace. One of the d i f f i c u l t i e s encountered i n producing  s t e e l samples from pure materials i s obtaining uniform carbon and a l l o y contents from accurately weighed materials.  Carbon  unites with the oxygen which i s always present i n varying quantities i n the melt, depending upon the purity of the materials used, the method of melting, time of melting, and degree of atmosphere protection given the melts.  Hence com-  plete carbon pickup i n the s t e e l from weighed amounts of pure graphite i s d i f f i c u l t unless complete atmosphere  protection  and similar melting conditions f o r each melt are maintained. Induction Furnace Reasonably uniform carbon analyses were obtained by melting i n the induction furnace, but several d i f f i c u l t i e s were experienced.  The temperature of the melt was hard to  control because the crucible d i d not heat uniformly nor to as high a temperature as the melt i t s e l f , and accurate measurement of the temperature of the s t e e l i t s e l f was not possible. The temperature of the melt was measured empirically by strap-  24. ping a platinum, platinum-rhodium thermocouple to the outside of the s i l i c a tube inside layers of asbestos i n s u l a t i o n ( f i g . 6), assuming 50 to 75 deg. P. l a g between the melt and the  thermocouple i t s e l f .  Melting proceeded u n t i l a Poten-  tiometer indicated the outside of the tube was 2900 deg. F. Time of melting f o r approximately 75 gram samples was one and one h a l f t o two minutes. ' Because the top of the c r u c i b l e was much colder than the melt i t s e l f considerable s t e e l froze during pouring. Hence the melts used f o r analyses were frozen i n the tube, which had to be broken to remove the ingot. A protective atmosphere could not be supplied without  a f f e c t i n g the temperature of the melt.  Five ingots gave  carbon analyses within a range of 0.05 percent carbon, but a 50 percent carbon loss was obtained when attempting to make a 0.50 carbon s t e e l . ferromolybdenum  Complete recovery of molybdenum from  and n i c k e l from n i c k e l shot was obtained.  Possible methods f o r producing uniform temperatures throughout the crucible and melt i n the Induction furnace are: (1) f i t a ready made crucible into a shaped graphite block, and (2) make a paste crucible of magnesia or z l r c o n l a inside a graphite mold and dry very slowly. Glo-bar Furnace The melt and crucible can be heated to a uniform controllable temperature i n the glo-bar furnace. Thus small  Old Method  t  CrucibU  B/och  Fig 7  for  fonovtd tcurinf  From  A/*t#MathoJ, fomortct  Crucible For  Pour  A/o+ my  Me/ting Crucible-and Carbon Block Support.  26. melts can be poured r e a d i l y without much cooling of the charge.  As the temperature required to melt the steels i s i n  the range 2950 to 2975 deg. P. i t i s necessary to have a crucible which has high thermal shock resistance and reasonably high strength at the high temperature.  Very few mater-  i a l s are s a t i s f a c t o r y or obtainable f o r use d i r e c t l y i n a b a s i c - l i n e d furnace.  High grade alundum c r u c i b l e s have  proved s a t i s f a c t o r y . To prevent oxidation and subsequent attack on the crucible l i n i n g ,  and also to reduce thermal and mechanical  shock, the alundum crucibles are placed i n a graphite block and covered with a graphite l i d .  In this way  reasonably  small carbon losses occur; sometimes, indeed, there i s a small carbon pickup from graphite powder f a l l i n g from the l i d . The method f i n a l l y adopted f o r making 100-gram s t e e l melts is as follows:(1)  Using Norway iron as the base material, holes  are d r i l l e d i n one end of small rods, into which are packed pure graphite powder and the required f e r r o a l l o y ( f e r r o molybdenum, ferrochromium, etc.) i n weighed amounts. (2)  The Norway iron bars are placed, packed-end  down, i n 35 ml porous alundum c r u c i b l e s , and any required amounts of n i c k e l added i n the form of shot. (3)  The c r u c i b l e i s placed i n a shaped graphite  block (fig.7) i n the furnace, covered with a graphite l i d ,  Sample ior  F19 8  Analysis  Carbon Mold For Castino Steel Melts.  28 and the furnace brought up to 2975 deg. P. (4)  A small amount of calcium-silicon i s added as  a deoxldizer and the furnace temperature which drops to 2925 is brought up to 2950 deg. P. (5)  The s t e e l i s poured into a heated graphite  mold ( f i g . 8) to make quarter-inch bars from which small specimens are cut after f u l l y annealing i n a s a l t bath f u r nace. Preliminary melts showed that, f o r Norway iron accurately packed with s u f f i c i e n t pure. graphite to make a 0.60 percent carbon s t e e l , the resultant product f o r three melts contained 0.57, 0.56, and 0.61 percent carbon respect i v e l y , which Indicates that reasonably close carbon content can be obtained by t h i s technique. Melts made without an addition of deoxldizer before pouring produced some hollow and some honeycombed bars.  Ex-  amination of every small specimen cut from eight deoxidized melts showed macro-cavities t o be p r a c t i c a l l y non-existent. Samples f o r analysis were machined from the top of the casting (see f i g . 6) which had been thoroughly cleaned and given the same softening treatment as the bars. The alundum crucibles used i n the glo-bar furnace withstand the high temperature reasonably w e l l a f t e r a thorough drying.  Although the crucibles are not attacked  chemically, i t was found that they are very susceptible to  29 thermal shock a f t e r one melt. Another type of crucible which may be used i n the glo-bar furnace to withstand the high temperatures and give atmosphere protection i s a paste of zlrconia or magnesia inside a graphite block.  This type has not been t r i e d .  Total recovery of added a l l o y i n g elements was tained i n a l l melts when n i c k e l shot, ferromolybdenum, ferrochromium were used.  oband  Very l i t t l e s i l i c o n was added to  the melts from the deoxldizer c a l c i u m - s i l i c o n .  EXPERIMENTAL TECHNIQUE IH ISOTHERMAL TRANSFORMATION Experimental technique i n producing Isothermal transformation diagrams was developed using a commercial a l l o y s t e e l NE8740.  The technique followed Is similar to the  methods developed and described i n detail.^ ^  wherein both  metallographic and dilatometric methods are used. Dilatometric specimens were cut l f Inch longitud i n a l l y from a 1% inch diameter bar of h o t - r o l l e d , a i r cooled NE8740 s t e e l . cut from 1/16  Metallographic specimens i  inch square were  inch transverse discs from the same bar.  Speci-  mens f o r microscopic work were wet-ground to a point approximately l/32 inch below the surface, near the center of a 1/16 by ij face, so that any decarburization was absent from the section examined, and so that the l o n g i t u d i n a l banding was  visible.  so. Since t h i s s t e e l showed marked banding, the c r i t e r ion of beginning of transformation was the f i r s t p r e c i p i t a t i n g p a r t i c l e s which could be detected v i s u a l l y at 1300 magnifications.  Because of the marked banding and the high magnlficao  tions required to detect each constituent, estimates of percentage transformation are d i f f i c u l t i f less than 10 percent i s transformed. A l l specimens were austenitized 15 minutes at 1550 deg. P. i n a neutral s a l t bath, producing average ASTM grain sizes 7 and 8.  The metallographic specimens were suspended  by 22 gauge chromel wire threaded through a small hole d r i l l e d i n the centre of a i - i n c h square face.  The d i l a t o -  meter specimens were handled with preheated chromel tongs. A l l transformations were c a r r i e d out i n s a l t baths controlled pyrometrlcally within plus or minus 5 deg. P.  Times were  measured from the instant the specimens entered the baths: less than two seconds was required to transfer the d i l a t o meter specimen from the a u s t e n i t i z i n g bath to the quench bath; less than o n e - f i f t h of a second was required to transfer a metallographic specimen from one bath to the next. Metallographic methods were found the more accurate. The dilatometer can only be used below 800 deg. P. and then only when the time to begin transformation i s long.  The Ms  l i n e was found by the method of Greninger and T r o i a n o , i n 3  which, b r i e f l y , specimens are quenched to form some martensite  doe  F/'Q9of  Seconds Beginning of I sot ft etm at Transformation Steel NB 6740 Austenitize of /SSoF  32 at a p a r t i c u l a r temperature  and held 10 seconds, raised above  the Ms Temperature to 700 deg. F. f o r 3 seconds to temper the martensite, and then quenched to room temperature.  The tem-  pered martensite etches dark whereas the untempered martensite does not. The beginning of transformation i s shown i n the form of an S-curve" ( f i g . 9) and data i s given i n Table 1. w  TABLE 1 Time to Begin Transformation f o r NE 8740 Steel  Temperature deg. P. 1250 1200 1100 1000 900 800 700  Time to Begin Transformation seconds  ~  10 8 6 5 5 5 10  TIME  go  MC HO  m  SECONDS  *»  Fig 10 Beginning of Transfor/notionof/Moy Contain I so thermally He/4 After Quenching Froml55  34. EXPERIMENTAL WORK The composition of Norway iron used as a base material of the alloys i s given i n Table 2 along with the analysis of the nickel-molybdenum series used f o r isothermal transformation diagrams.  TABLE 2 Analysis of Alloys Investigated and Norway Iron Alloy No.  Ae^ Chemical composition percent Temperature C Mn Si Nl Mo deg. F.  Norway iron  0.02 <0.001  0.004  A  0.55  0.005  AA1  0.55  0.  0.02  AA2  0.50  0.  AA3  0.59  0.  Grain Size ASTM  0.35  1333  7-8  2.13  0.34  1275  8  0.01  3.61  0.35  1210  9  0.02  5.39  0.35  1160  7-9  The method used f o r isothermal transformation of the a l l o y steels was the metallographic technique developed Bain^.  by  To conserve specimens i n f i n d i n g the beginning of  transformation times used for preliminary surveys were f i v e or ten second i n t e r v a l s . transformation was  The exact time of beginning of  found by a f i n a l survey at times between  the l a s t one showing no transformation and the f i r s t one showing transformation has begun.  1406  5£C0W5 FifH- Beqinninoof Trans formaf/on of Alloy Containing 2'13%11'i Isothermal/*} Held After Quenching  From  1550  F-  36. Specimens f o r isothermal transformation were cut approximately one-twelfth of an inch thick from the cast bars which had been annealed. center of each specimen  Small holes were d r i l l e d i n the through which 22 gauge chromel wire  was threaded to serve as a handle f o r t r a n s f e r r i n g the specimen from one bath to the next.  Three minutes at a u s t e n i t l z i n g  temperature of 1550 deg. P. were required to completely aust e n i t i z e the specimen, and hence a l l specimens isothermally treated were heated s i x minutes at 1550 deg. P. i n a neutral s a l t bath.  Transformation temperatures used were 1200, 1100,  1000, 930, 880, 800, 750, 700, and 650, a l l i n deg. P.  Speci-  mens were quenched f o r a measured time at one of the above temperatures i n a s a l t bath, followed by an immediate quench to room temperature i n an 8 percent brine s o l u t i o n . Th© f l a t faces of the transformed specimens were smoothed on No. 2 emery paper, the specimens mounted i n transoptic, polished, etched, examined a t 1800 magnifications, and a suitable series of specimens photographed to i l l u s t r a t e beginning of decomposition where one product i s formed, and to i l l u s t r a t e progress i n transformation where more than one product Is formed.  14C0  1300  W j .  !  g  S  '*  **  TIME  SO  Fi'o/2 Bo$ir*nit%j of Transformation  Isothermal/^  Held After  HQ  SCO  SOO /OOO XOOO 9OO0  SECONDS  of Alley Contmininq 3-6$  Quench iny Frem  ittO  F  38.  EXPERIMENTAL RESULTS The beginning of isothermal transformation of austenite f o r the series of 0.55 carbon, 0.35 molybdenum, 0, 2.13, 3.69, and 5.31 percent n i c k e l austenitized s i x minutes at 1550 deg. P. i s shown i n f i g . 10 to 13. The p e a r l i t e reaction i s delayed much more than can be explained by the lowering of the equilibrium temperature. n i c k e l produces a bay i n the curve at 880 deg. P. formation i n the Intermediate  Increasing Ferrite  temperature range 880 to 1000  deg. F. becomes Increasingly more a c i c u l a r , and during growth of the acicular f e r r i t e carbides are rejected e i t h e r from the f e r r i t e or at the f e r r l t e - a u s t e n i t e Interface, followed at l a t e r times by growth of the carbide phase, at f i r s t between groups of a c i c u l a r f e r r i t e and eventually into the untransformed austenite surrounding the f e r r i t e . The results showing the e f f e c t of n i c k e l on the isothermal transformation products of austenite are shown as photomicrographs i n f i g . 14 to 101.  1400  1300  ****  1  2  r  To90  TIME Figl3  So  too too  SOP 'too emsooo  tooao  SECONDS  Beginning of Transformation of Alloy Contain inj 53l%Hi* I sot normally Held fiftor Quench in $ From IS 50 F.  Alloy A  (0 percent nickel) A very f i n e dendritic pattern i s produced on cast-  ing a l l o y A ( f i g . 14) which i s a regular pattern of Widraanstaaten f e r r i t e and fine p e a r l i t e ( f i g . 15).  Annealing by  heating f o r 30 minutes at 1550 deg. P. and cooling slowly i n f i v e hours to room temperature produces a homogeneous f i n e grained structure ( f i g . 16) of f e r r i t e and coarse lamellar p e a r l i t e ( f i g . 17), the structure which is used f o r isothermal work. The transformation product between 650 and 880 deg. P. i s feathery lower bainite ( f i g . 18 to 25).  At 930 deg. P.  upper bainite grows nodularly ( f i g . 26), the structure of which i s not resolved at 1800 magnifications.  At 1000 deg.P.  two products form and grow simultaneously ( f i g . 27, 28 and 29), the l i g h t - e t c h i n g phase being a c i c u l a r f e r r i t e , and a darketching phase nodular i n formation and very rapid-growing. At 1100 deg. P. a c i c u l a r f e r r i t e forms at two seconds ( f i g . 30) and a dark-etching phase begins to form at f i v e seconds near adjoining groups of f e r r i t e fingers ( f i g . 31).  At 1200 deg.  P. F e r r i t e begins at f i v e seconds ( f i g . 32) and has an acicular structure which becomes more rounded by the time p e a r l i t e forms at 15 seconds ( f i g . 33).  The time of beginning  of reaction between 800 and 1000 deg. F. i s undetermined because i t i s below the l i m i t of accurately measuring time at temperature.  F i g . 14 A l l o y A as cast X100 Fine dendritic pattern of f e r rite(white) and p e a r l i t e ( b l a c k ) .  F i g . 15 A l l o y A as cast X850 Widmanstaaten ferrite(white)  F i g . 16 A l l o y A annealed X100 Fine grain structure of f e r r i t e (white) and coarse pearlite(dark)  F i g . 17 A l l o y A anneal* Lace network of coarse and f i n e p e a r l i t e plus f e r r i t e . ~  l  —  4 ^  F i g . 18 A l l o y A 3 sec. 650 deg. F. shows bainite (dark specks) and martensite X850.  _  ,.:.r..,:sx.  • ¥  F i g . 19 A l l o y A 5 sec. at 650 deg. F. shows growth of bainite. X850  «  • •** Twain *  /  ,*  N"  F i g . 20 A l l o y A 2 sec. at 700 deg. F. Bainite(black) begins. Matrix martensite. X850  F i g . 21 A l l o y A 3 sec. at 700 deg. F. Bainite grows. X 850  ,j  i  ,  '»  u V>  0  ' it !' '  rM  *  • *  K.  *  F i g . 22 A l l o y A 1 sec. at 750 deg. F. Bainite forms. X1000  F i g . 23 A l l o y A Bainite after 3 s e c at 750 deg. F. X 850  --—r-w  v.  •v.  ;  1  1  ,•  •& "  1  9  ..•  *  .  »,  .  4  .  F i g . 24 A l l o y A Bainite formed at 1 sec. 800 deg. F. X850  F i g . 25 A l l o y A transformed 1 see. at 880 deg. F. X1000  F i g . 26 A l l o y A transformed 1 seo. at 930 deg. F. t o upper bainite. X1000  F i g . 27 Alloy A transformed 1 sec. at 1000 deg. F. Acicular ferr i t e ( l i g h t ) and dark phase. X850  I. F i g . 28 Alloy A transformed 2 sec. at 1000 deg. F. Darketching phase grows. X850  F i g . 29 A l l o y A transformed 5 sec. at 1000 deg. F. X850  » W W .•SET-?*' * ^ 3  F i g . 30 Alloy A transformed 3 seo. at 1100 deg. F. to acicular proeutectoid f e r r i t e . X850  I  F i g . 31 A l l o y A transformed 5 sec. at 1100 deg. F. p e a r l i t e begins t o form. X850  F i g . 32 A l l o y A transformed 5 sec. at 1200 deg. F. Proeutectoid f e r r i t e begins to form. X850  F i g . 33 A l l o y A transformed 15 sec. at 1200 deg. F. P e a r l i t e begins t o form(center). X850  A l l o y AAi  (2.13 percent nickel) A very fine dendritic pattern i s developed i n the  cast ingots ( f i g . 34) which i s an i n t r i c a t e pattern of f e r r i t e fingers and p e a r l i t e ( f i g . 35). Annealing the a l l o y removes the d e n t r i t i c pattern giving a f i n e grained (fig.36), homogeneous mixture of f e r r i t e and f i n e lamellar p e a r l i t e ( f i g . 37) which i s the structure previous to isothermal heat treatment. Between 650 and 880 deg. P. feathery bainite forms ( f i g . 38 to 45). At 930 deg. P. a nodular dark-etching product forms ( f i g . 46) and i s probably upper b a i n i t e .  At  1000 deg. P. two products appear simultaneously ( f i g . 47), acicular f e r r i t e and a nodular, dark-etching constituent which grows more r a p i d l y than f e r r i t e ( f i g . 48). At 1100 deg. P. f e r r i t e precipitates a t 5 seconds ( f i g . 49) and grows alone u n t i l 400 seconds when p e a r l i t e appears  ( f i g . 50). At 1200  deg. P. f e r r i t e begins to form at 20 seconds ( f i g . 52) but p e a r l i t e was not found up to 20,000 seconds.  )  F i g . 34 A l l o y AA-^ as cast Ferrite(white),pe a r l i t e ( b l a c k ) , fine dendrites. X100 n  r  F i g . 36 A l l o y AA^ as cast shows feathery i n t e r l a c i n g Widmanstaaten f e r r i t e X850.  .  f * -«  F i g . 36 A l l o y AA^ Annealed Fine-grained structure of f e r r i t e (white) and p e a r l i t e (dark). X100  F i g . 37 A l l o y AA^ annealed Fine pearlite and f e r r i t e X850  T ' 3 - 9  I  \  \ • I  Fig.38 A l l o y AA^ transformed 15 sec. at 650 deg. F. to show begin of "bainite X1200  F i g . 39 A l l o y AAj transformed 17 sec. at 650 deg. F. X850  AT _ F i g . 40 A l l o y AA^ transformed 10 s e c at 700 deg. F. Bainite begins t o form. X850  F i g . 41 A l l o y AAi transformed 12 sec. at 700 deg. F. X850  F i g . 42 A l l o y AA^ transformed 4 sec. at 750 deg. F. Bainite begins t o foim. X850  F i g . 43 A l l o y AA-^ transformed 3 sec. at 800 deg. F. Bainite begins to form. X850  ' ' J  * * * * * *  3 »  *  I  .  F i g . 44 Alloy AA^ transformed 2 sec. at 880 deg. F. Bainite begins to form. X850  F i g . 45 A l l o y AA^ transformed 3 sec. at 880 deg. F. X850.  F i g . 46 A l l o y AA^ transformed 3 sec. at 930 deg. F. X850 Upper bainite begins t o form  F i g . 47 A l l e y AA]_ transformed 3 sec. at 1000 deg. F. Acicular f e r r i t e and a dark-etching phase  F i g . 50 A l l o y AAj transformed 400 sec. at 110O deg. F. Pearlite begins t o f o m . X850  F i g . 51 Alloy AAi transformed 500 sec. at 1100 deg. F. X850  49  <  m  T F i g . 52 A l l o y AA^ transformed 20 sec. at 1200 deg. F. F e r r i t e begins t o form. X850  F i g . 53 A l l o y AA^ transformed 40 sec. at 1200 deg. F. X850  F i g . 54 A l l o y AA^ transformed 5000 sec. at 1200 deg. F. X850  A l l o y AAg  (3.61 percent n i c k e l ) A very fine d e n d r i t i c pattern i s produced on  casting ( f i g . 55) which i s a very i n t r i c a t e pattern of acicular f e r r i t e and b a i n i t e ( f i g . 56). Annealing the a l l o y removes the dendrite pattern, producing a uniform fine-grained structure ( f i g . 57). Several products form during annealing, among which are proeutectoid f e r r i t e , fine p e a r l i t e , acicular f e r r i t e , and a secondary carbide produced at the same temperature range as the a c i c u l a r f e r r i t e and b a i n i t e ( f i g . 58). Feathery b a i n i t e forms between 650 and 800 deg. F. ( f i g . 59 to 62).  At 880 deg. F. the f i r s t product a f t e r 9  seconds appears to be a c i c u l a r f e r r i t e ( f i g . 63).  Further  progress of transformation at 880 deg. F. shows two products, the feathery or acicular product being f e r r i t e with r e j e c t i o n of a carbide between 20 and 25 seconds ( f i g . 64 and 65), the carbide then growing separately within areas of f e r r i t e ( f i g . 66).  Similar growth i s found at 930 deg. F.; a f t e r 10  seconds a c i c u l a r f e r r i t e p r e c i p i t a t e s ( f i g . 67) followed at 15 seconds by r e j e c t i o n of carbide ( f i g . 68) and growth of the carbide phase separately within groups of a c i c u l a r f e r r i t e at 20 seconds ( f i g . 69) and eventually away from f e r r i t e groups by 30 seconds ( f i g . 70). At 1000 deg. F. f e r r i t e p r e c i p i t a t e s a f t e r 20 seconds ( f i g . 71) and grows  u n t i l at 400 seconds carbide precipitates  ( f i g . 72) and grows  as a separate phase between 500 and 1000 seconds ( f i g . 75 and 74).  At 1100 deg. P. f e r r i t e forms at 25 seconds ( f i g . 75)  and grows u n t i l 2000 seconds when f i n e lamellar p e a r l i t e begins to grow ( f i g . 76). At 1200 deg. P. f e r r i t e p r e c i p i tates at 500 seconds ( f i g . 77) and grows rounded and agglomerated by 10000 seconds ( f i g . 78). P e a r l i t e was not found up to 20000 seconds.  F i g . 55 Alloy AAg as cast showing extremely fine dendrites. X100  F i g . 56 A l l o y AAg as cast showing acicular f e r r i t e and bainite. X850  F i g . 57 Alloy AA annealed shows extremely f i n e grain structure. X100  F i g . 58 A l l o y AAg annealed X850 Proeutectoid and a c i c u l a r f e r r i t e P e a r l i t e and secondary carbide.  2  Si  V  Vt i  '•  4/  F i g . 59 A l l o y AAg transformed 30 seo. at 650 deg. F. Bainite begins t o form. X850  F i g . 60 Alloy AAg transformed 25 sec. at 700 deg. F. Bainite begins to form. X850  -  V ,  F i g . 61 A l l o y AA^ transformed 9 sec. at 750 deg. F. Bainite begins t o form. X850 •  F i g . 62 Alloy AAg transformed 7 sec. at 800 deg. F. Bainite begins to form. X850  •7  »  9 •  f  2«r  F i g . 63 Alloy AAg transformed 9 sec. at 880 deg. F. Acicular f e r r i t e begins. X850 P i c r a l .  4  F i g . 64 A l l o y AA> transformed 20 sec. at 880 deg. F. aoioular f e r r i t e grows. X850  •*  X  *• 1 J  S  50 F i g . 65 A l l o y AA£ transformed 25 seconds at 880 deg. F. shows carbide rejection from f e r r i t e . X850  F i g . 66 A l l o y AA£ transformed 40 s e c at 880 deg. F. shows growth of the carbide as a separate phase. X850  F i g . 67 A l l o y AAg transformed 10 sec. at 930 deg. F. F e r r i t e begins t o form. X850  F i g . 68 A l l o y AAg transformed 15 sec. at 930 deg. F. Carbide rejected from a c i c u l a r f e r r i t e . X850.  i  F i g . 69 Alloy AAg transformed 20 sec. at 930 deg. P. Carbide begins t o grow. X850  F i g . 70 Alloy AA transformed 30 sec. at 930 deg. F. Carbide grows r a p i d l y . X850 2  7  F i g . 71 A l l o y AA transformed 10 sec. at 1000 dSg. F. F e r r i t e begins to form. X850 ?  F i g . 72 A l l o y AA, transformed 400 sec. at 1000 fleg. F. Carbide begins to p r e c i p i t a t e . X850  F i g . 73 A l l o y AA» transformed 500 sec. at 1000 deg. F. X850 Carbide begins t o grow.  F i g . 74 A l l o y AAg transformed 1000 sec. at 1000 deg. F. X850 Further carbide growth.  F i g . 75 A l l o y AA« transformed 20 sec. at 1100 deg. F. F e r r i t e begins t o form. X850  F i g . 76 A l l o y AAg transformed 2000 sec. at 1100 deg. F. X850 Fine pearlite begins t o form.  7? * \ ,  J  F i g . 77 A l l o y AAg transformed 500 sec. at 1200 deg. F. Ferrite begins to form. X850  F i g . 78 A l l o y AA transformed 10000 seo. at 1200 deg. F. X850 2  A l l o y AA3  (5.31 percent n i c k e l ) The structure developed on casting i s a very f i n e  dendritic pattern ( f i g . 79) of feathery f e r r i t e along Widmanstaaten planes  ( f i g . 80) bainite and martensite.  Annealing the cast ingot does not remove e n t i r e l y the dend r i t i c pattern ( f i g . 81) but leaves a very f i n e network of Widmanstaaten f e r r i t e , a c i c u l a r f e r r i t e , carbide rejected and grown from a c i c u l a r f e r r i t e , b a i n i t e and martensite ( f i g . 82).  Between 650 and 800 deg. P. feathery b a i n i t e  forms and grows ( f i g . 83 to 87).  At 880 deg. P. F e r r i t e  p a r t i c l e s p r e c i p i t a t e at 25 seconds ( f i g . 88) followed by carbide r e j e c t i o n at 30 seconds ( f i g . 89) as the f e r r i t e grows a c i c u l a r l y , and the carbide phase begins to grow at 60 seconds ( f i g . 90 and 91).  At 930 deg. F. a c i c u l a r f e r r i t e  forms a f t e r 15 seconds ( f i g . 92), followed by carbide r e j e c t ion at 30 seconds on growth of the f e r r i t e ( f i g . 93 and 94), and the carbide phase begins to grow at 80 seconds ( f i g . 95) within groups of f e r r i t e and extends beyond f e r r i t e groups by 120 seconds ( f i g . 96).  At 1000 deg. F. f e r r i t e fingers form  at 30 seconds ( f i g . 97) and grows rounded ( f i g . 98 to 102). P e a r l i t e was not found at 20000 seconds, at which temperature 1  f e r r i t e forms as a line marking grain boundaries as w e l l as the large p a r t i c l e s ( f i g . 103).  At 1100 deg. P. f e r r i t e be-  gins after 100 seconds ( f i g . 104), and grows rounded by 10000 seconds ( f i g . 105). The p e a r l i t e reaction does not appear at 20000 seconds.  F i g . 79 A l l o y AA- cast and drawn 1100 deg. F. X100 Fine dendritic pattern  F i g . 80 A l l o y AAg cast and drawn 1100 deg. F. X850 F e r r i t e and tempered bainite & martensite.  f m  .  s  ?  m  m  F i g . 81 A l l o y AA annealed X100 Dendritic pattern not e n t i r e l y removed• 3  F i g . 82 A l l o y AA3 annealed. X850 F e r r i t e , secondary carbide, bainite and martensite.  Ik  F i g . 83 A l l o y AAg transformed 60 sec. at 650 deg. F. Bainite begins t o form. X850  F i g . 84 Alloy AA3 transformed 30 sec. at 700 deg. F. Bainite begins t o form. X850  1V  -»-  F i g . 85 A l l o y AA» transformed 25 seo. at 750 deg. F. Bainite begins to form. X850  F i g . 86 A l l o y AA- transformed 14 seo. at 800 deg. F. Bainite begins t o foim. X850  » R  *  '• f •  F i g . 87 A l l o y AAg transformed 20 sec. at 800 deg. F. Bainite begins to form. X850  * 0£  F i g . 88 A l l o y AA3 transformed 25 sec. at 880 deg. F. F e r r i t e begins t o Form. X1000  MS- <  F i g . 89 A l l o y AAg transformed 30 sec. at 880 deg. F. X850 F e r r i t e grows a c i c u l a r l y , rejecting carbide.  F i g . 90 A l l o y AA3 transformed 60 sec. at 880 deg. F. X850 Carbide begins t o grow.  F i g . 91 A l l o y AAg transformed 120 sec. at 880 deg. F. X850 Further growth of carbide.  is* C  F i g . 92 Alloy AAg transformed 15 sec. a t 930 deg. F. F e r r i t e begins t o form. X850  '  F i g . 93 A l l o y AAg transformed 30 sec. at 930 deg. F. X850 Carbide rejected as f e r r i t e grows  F i g . 94 A l l o y AAg transformed 40 sec. at 930 deg. F. X1000 Further carbide r e j e c t i o n .  F i g . 95 A l l o y AAg transformed 80 sec. at 930 deg. F. X850 Carbide phase begins t o grow.  F i g . 96 A l l o y AAg transformed 120 sec. at 930 deg. F. X850 Carbide growth much greater.  •  F i g . 97 A l l o y AA transformed 30 sec. at 1000 deg. F. F e r r i t e begins t o form. X850 S  F i g . 98 A l l o y AA* transformed 150 sec. at 1000 deg. F. F e r r i t e grows rounded. X850  •  F i g . 99 A l l o y AA transformed 300 sec. at 1000 deg. F. X850 F e r r i t e grows further. 3  F i g . 100 A l l o y AA? transformed 500 seo. at 1000 deg. F.X850  >  I  L  4 * ..  •  *; Pi  5  F i g . 101 A l l o y AAg transformed 1000 sec. at 1000 30 deg. dee;. F. X850  3 V: < \  £ n  F i g . 102J Alloy AA transformed 2000 sec. a t 1000 deg. F. X850 3  61.  F i g . 103 A l l o y AAg transformed 20000 sec. at 1000 deg. F. showing f e r r i t e outlining grain boundaries. X850  F i g . 104 A l l o y AAg transformed 100 sec. at 1100 deg. F. F e r r i t e begins to form. X850  F i g . 105 A l l o y AAg transformed 10000 sec. at 1100 deg. F. F e r r i t e grown similar t o that at 1000 deg. F. X850  62 DISCUSSION OF RESULTS ' Several Investigators have reported that n i c k e l s h i f t s the whole isothermal transformation curve to the  right,  and that a l l portions of the curve are s h i f t e d similar amounts.  The present investigation has found ( f i g . 10 to  13)  that the p e a r l i t e curve i s pushed to the right much more than any other part of the isothermal transformation curve with increasing n i c k e l content.  The  e f f e c t on the p e a r l i t e r e -  action cannot therefore be d i r e c t l y related to the increase i n n i c k e l content.  To f i n d the d i r e c t e f f e c t of n i c k e l on  the p e a r l i t e reaction a series of pure iron-nickel-carbon alloys s h a l l be investigated.  If the p e a r l i t e reaction  be explained by the lowering of the eutectoid  temperature the  curves of t h i s reaction drawn with a common Ae would be i d e n t i c a l .  temperature  It can therefore only be assumed that  n i c k e l i n t e n s i f i e s greatly the e f f e c t of molybdenum on pearlite  the  reaction. By extrapolation  and Ham^  could  of data given by Wells and Mehl^g  the approximate values of d i f f u s i o n c o e f f i c i e n t s of  n i c k e l and molybdenum at 1000  deg. P. i n austenite containing  0.6  14  percent carbon are 9 x 1 0 ~  respectively,  and 8 x IO - - cm -  compared with 8 x 10"  The r a t i o of d i f f u s i o n rates i s PJ2 DNj  9  cm  2  = 9.  1  3  2  per second  per second f o r carbon 15 The  r a t i o at  1550  T  deg. P. i s 4 to 1; thus, i n the range of temperature which  63. p e a r l i t e forms n i c k e l diffuses more slowly than does molybdenum.  Increasing the n i c k e l content should increase the  preference i n forming f e r r i t e n u c l e i , because n i c k e l i s an element p a r t i t i o n e d almost e n t i r e l y i n f e r r i t e .  The delay i n  forming f e r r i t e at equivalent temperatures below the e q u i l i brium (the. gamma to gamma plus alpha equilibrium i s lowered proportionally to the eutectoid lowering with increasing i  nickelg^) may be explained on a basis of lowered d i f f u s i o n . But the delay i n the p e a r l i t e reaction may not e n t i r e l y be the r e s u l t of lowered d i f f u s i o n . The increase i n the f e r r i t e nucleation rate by n i c k e l and the increase i n the carbide nucleation rate by molybdenum are two opposing factors which depend on free energy and d i f f u s i o n rates before growth of n u c l e i can occur. Because the p e a r l i t e reaction i n the higher n i c k e l a l l o y s tested isothermally did not appear within the time l i m i t s i n vestigated (assuming the reaction does appear at longer times according to Bain's equilibrium phase diagram^) even a f t e r the f e r r i t e has apparently grown to i t s f u l l amount, some other cause besides d i f f u s i o n i s necessary to explain the delay.  This suggests two possible hypotheses; (1) that  nucleation s i t e s f o r carbides are increasingly more d i f f i c u l t to f i n d as the austenite n i c k e l content i s increased, and that these carbide nucleation s i t e s cannot exist u n t i l f e r r i t e  64,. transformation  and further n i c k e l d i f f u s i o n from austenite  to  f e r r i t e lowers the n i c k e l content of the remaining austenite, or (2) that n i c k e l increases the tendency f o r carbide n u c l e i to be less than the c r i t i c a l size f o r growth according  to free  energy r e l a t i o n s h i p s . The accepted theory of nucleation and growth of nuclei- - . . ^ - . ^ ^ 5 , y ,10,11,lo 6i Q  r  o r 7 f  i n  ,  s o l i d - s o l i d reactions  (based on  analogy of g a s - l i q u i d - s o l i d reactions) i s as follows.  During  the random migrations of atoms i n austenite, volumes are established from time to time of the required l a t t i c e f o r a p r e c i p i t a t i n g phase; the number of these embryo l a t t i c e s or potential nuclei depends on d i f f u s i o n r e l a t i o n s h i p s ; and  the  size of the p o t e n t i a l nuclei i s s t a t i s t i c a l l y d i s t r i b u t e d . Free energy relationships determine the minimum s i z e p o t e n t i a l nucleus which has a surface energy equivalent  to the difference  i n free energy between austenite and a phase of the p o t e n t i a l nucleus l a t t i c e .  I f this condition Is met by a p o t e n t i a l  nucleus i t can remain i n s i t u as a stable nucleus, and growth of t h i s nucleus into a phase of the new  l a t t i c e can occur as  d i f f u s i o n i n austenite brings other atoms appearing i n the phase to l a t t i c e positions adjoining the nucleus. condition i s not met  new  If this  the embryo l a t t i c e i s broken up by con-  tinued d i f f u s i o n of the atoms.  But the accepted theory has  been questioned to explain some discrepancies, the argument being that thermodynamics does not t e l l us exactly what happens,  65. but what Is most l i k e l y to happen.  Hultgren  has suggested  g  growth from an embryo l a t t i c e smaller than the  accepted  minimum, or, f o r isothermal transformation, growth from zero time. This leads to the hypothesis discussed  earlier  that n i c k e l I n t e n s i f i e s formation of smaller embryo l a t t i c e s or "quasi-nuclei" of carbides, and that growth of these smaller n u c l e i must necessarily take longer than growth from much larger n u c l e i , which results i n transformation time to v i s i b l e amounts being extremely long; or that the "quasin u c l e i " themselves have to diffuse before growth can occur, and d i f f u s i o n rate of a "quasi-nucleus" i s smaller than that of n i c k e l and molybdenum.  This hypothesis may  explain some  of the f a c t s better than other hypotheses and can only be proven by Indirect means. The e f f e c t of n i c k e l i n Increasing both the amount and time of formation of a c i c u l a r f e r r i t e at  Intermediate  temperatures can r e a d i l y be explained by the increased tendency toward f e r r i t e nucleation and the low d i f f u s i o n rate of n i c k e l causing slower growth of f e r r i t e n u c l e i .  Lower  d i f f u s i o n rates of a l l elements at these temperatures and the increased nucleation rate of f e r r i t e results i n the formation of a high carbon f e r r i t e .  As the a c i c u l a r f e r r i t e grows  further, d i f f u s i o n of carbon and carbide-forming  elements  causes p r e c i p i t a t i o n of carbides, both within the f e r r i t e  66 and at the f e r r i t e - a u s t e n i t e i n t e r f a c e . carbide grows as a separate phase.  Eventually, the  This type of p r e c i p i t a -  t i o n has been discussed before i n the l i t e r a t u r e .  Davenport  c a l l s a c i c u l a r f e r r i t e the "X" constituent and notes the secondary carbide reaction i n low and medium a l l o y s t e e l s . This type of transformation was also notedgg i n transformations of low and medium carbon N E s t e e l s , and i n a l l a l l o y s discussed  the p e a r l i t e reaction was delayed greatly. The bay which appears i n the curve has not been  noted before f o r n i c k e l s t e e l s .  The lower b a i n i t e reaction  i s evidently separate from the f e r r i t e reaction.  The upper  bainite reaction has been displaced by the a c i c u l a r f e r r i t e reaction, and the bay i n the curve between the a c i c u l a r f e r r i t e and lower bainite reactions i l l u s t r a t e s the separat i o n i n maxima of the two reactions.  67 CONCLUSIONS (1)  Nickel s h i f t s the isothermal transformation  curve generally to the r i g h t i n a 0,55 carbon, 0.35 molybdenum a l l o y of i r o n . (2)  Nickel delays the p e a r l i t e reaction more than  can be explained by the lowering of the eutectoid temperature. To f i n d the actual e f f e c t of n i c k e l and molybdenum, as' single alloying elements and together, isothermal transformation data w i l l have to be obtained f o r pure iron-nickel-carbon alloys and pure Iron-molybdenum-carbon a l l o y s , and the results of these investigations compared with those of the present investigation. (3)  Nickel produces a tendency toward formation of  high carbon a c i c u l a r f e r r i t e at intermediate  temperatures,  accompanied on further growth of the f e r r i t e by carbide prec i p i t a t i o n and agglomeration, and by eventual growth of the carbide as a separate phase formed d i r e c t l y from austenite. (4)  The upper bainite reaction i s suppressed as  n i c k e l content Is increased, and i s displaced by the acicular f e r r i t e reaction. (5)  The lower bainite reaction as represented on  the curve becomes separated from the rest of the curve by the appearance of a bay between the acicular f e r r i t e and bainite reactions, a f a c t which has never been reported i n  68.  e a r l i e r investigations of alloys containing n i c k e l . (6)  The upper temperature of formation of bainite  i s not lowered as r a p i d l y as the eutectoid increasing n i c k e l .  temperature with  69.  ACKNOWLEDGMENTS  This research has been conducted under grants of the University of B r i t i s h Columbia Alloys Research fund and grants of the National Research Council.  This thesis  i s part o f an extended study under National Research Council  sponsorship.  The author wishes to thank Professor P. A. Forward, head of the Department, and Associate Professor W. M. Armstrong, Department of Mining and Metallurgy, f o r many suggestions and c r i t i c i s m s of t h i s research.  70 BIBLIOGRAPHY 1.  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