200903532 九、發明說明 【發明所屬之技術領域】 本發明係關於R-Τ-Β系合金' R-Τ-Β系稀土類永久磁 鐵用微細粉末、R-T-B系稀土類永久磁鐵，特別是可獲得 保磁力優異之R-Τ-Β系稀土類永久磁鐵之R-Τ-Β系合金及 R-Τ-Β系稀土類永久磁鐵用微細粉末。 【先前技術】 R-Τ-Β系磁鐵由於其高特性而常使用於硬碟（HD)、 磁氣共振影像法（MRI )、各種馬達等。近年來由於對R_ T-B系磁鐵耐熱性的提升及節約能源的要求提高，因而包 含汽車之馬達用途比率提昇。 R-Τ-Β系磁鐵，由於其主成分爲Nd、Fe、B，故統稱 爲Nd-Fe-B系或R-Τ-Β系磁鐵。R-Τ-Β系磁鐵之R係Nd 之一部分用Pr、Dy、Tb等其他稀土類元素置換者等。τ 爲Fe之一部分用Co、Ni等其他過渡金屬置換者。B爲硼 ，其一部分可用C或N置換。 R - T - B系磁鐵之R - T - B系合金係有助於磁化作用之磁 性相的R2TMB相所成之主相’以及非磁性稀土元素濃縮 後低溶點之昌R相並存之合金。由於R-Τ-Β系合金因爲活 性金屬，故一般可於真空或惰性氣體中融化或鑄造。又， 有關經鑄造之R -τ B系合金根據粉末冶金法製作燒結磁鐵 方面，一般係將合金塊粉碎至平均粒徑5ym左右（d50: 以激光繞射式粒度分佈計測量）而形成合金粉末後，在磁 -4 - 200903532 場中擠壓成形’於燒結爐以約1 0 0 0至11 0 0 °C之高溫燒結 ，之後視需要進行熱處理、機械加工’更爲了提高耐蝕性 而施予鍍覆，作成燒結磁鐵。 於R-T-B系燒結磁鐵中，富R相係擔當如下述之重要 角色’· 1 )熔點低，燒結時成爲液相，磁鐵高密度化，因此有助 於提高磁性。 2 )使粒界之凹凸消失，減少逆磁區之成核點（nucleation s i t e )而提高保磁力。 3 )使主相磁性絕緣而增加保磁力。 因此，若成形之磁鐵中富R相之分散狀態不佳時則會 招致局部性燒結不良、磁性降低，因而富R相平均分散於 成形之磁鐵中甚爲重要。 R-T-B系燒結磁鐵之富R相之分佈受到原料R-T-B系 合金組織的影響很大。 又，R-T-B系合金鑄造中衍生之另一問題係經鑄造之 合金中生成a -Fe。由於α -Fe具有變形力，未粉碎而殘留 於粉碎機中，故不僅使粉碎合金時之粉碎效率降低，亦影 響到粉碎前後之組成變動、粒度分佈。再者，若a -Fe燒 結後仍殘留過磁鐵中，則會造成磁鐵之磁特性減弱。所以 ’以往之合金係視需要在高溫下進行長時間之均質化處理 ’以消除a - F e。但因α _ ρ e係以包晶核存在，故其消除必 須爲長時間的固相擴散，若爲數cm厚度之鑄錠（ingot ) 而稀土類量爲33%以下則事實上不可能消除α -Fe。 200903532 爲了解決此R-Τ-Β系合金中產出α -Fe之問題，發展 出以急速冷卻速度鑄造合金塊之薄帶連鑄（stripcast)法 (簡稱SC法），並實用化。SC法係藉由於內部經水冷的 銅滾筒上流入熔液鑄造〇·】至1mm左右之薄片，而將合金 急速凝固之方法。就SC法而言，由於將熔液（molten metal)過度冷卻至主相R2T14B相的生成溫度以下，故可 直接從合金熔液生成R2T14B相，而可抑制a -Fe的澱積。 再者，藉由進行S C法，因其合金的結晶組織微細化，故 可生成具有富R相係經微細分散之組織之合金。富R相在 氫氣氛圍中與氫反應，變成膨脹且脆的氫化物。若利用此 性質，則可導入對應於富R相分散程度之微細裂縫。若經 由氫化步驟後再進行微細粉碎，則由於利用氫化所生成的 多量微細裂縫使合金遭破壞而粉碎性變得極佳。如此，用 SC法鑄造之合金，由於其內部富R相呈微細地分散，故 粉碎、燒結後之磁鐵中之富R相之分散性亦良好，而成功 提高了磁鐵之磁特性（例如參照專利文獻1)。 又，以SC法鑄造之合金薄片，其組織之均質性佳。 組織之均質性可用結晶粒徑或富R相之分散狀態加以比較 。用SC法製得之合金薄片，在合金薄片鑄造用滾筒側（ 以下亦稱爲鑄型面側）雖會產生激冷晶（c h i 11 c r y s t a 1 )， 但可獲得因整體急速冷卻凝固所造成的適度微細而均質之 組織。 如上所述，用SC法鑄得之R-T-B系合金，由於富R 相微細分散，a -Fe之生成亦受抑制，因而具有製作燒結 -6 - 200903532 磁鐵用之優異組織。 就磁特性，特別是保磁力與磁鐵微細構造中的元素分 佈之關係而言’對有助於保磁力提高的D y之分佈影響很 大。例如：已有報告指出若Dy分佈於晶界相附近時則保 磁力提高（例如參照專利文獻2)。 更詳細而言’亦有報告指出若該等Dy存在於主相時 保磁力會提高（例如參照專利文獻3、非專利文獻1 )。 又’因爲磁鐵的特性與合金製造方法有一定的關聯性 ’故隨著磁鐵特性提昇，其合金之製造方法亦進步。例如 :已知的控制微細構造之方法（例如參照專利文獻4 )、 將鑄造滾筒的表面狀態加工至預定的粗度而控制微細構造 之方法（例如參照專利文獻5、專利文獻6 )。 〔專利文獻1〕日本特開平5 - 2 2 2 4 8 8號專利公報 〔專利文獻2〕日本特開平5 _ 2 1 2 1 9號專利公報 〔專利文獻 3〕W02003/001541 〔專利文獻 4〕W02005/031023 〔專利文獻5〕日本特開2 〇 0 3 _丨8 8 〇 〇 6號專利公報 〔專利文獻6〕日本特開2〇〇4_4329ι號專利公報 〔非專利文獻1〕富澤浩之、粉體及粉末冶金、 2005年3月第52卷第3號第158至163頁。 【發明內容】 發明欲解決之問題 然而’近年來追求更高性能之R_T_B系稀土類永久磁 200903532 鐵，需要更進一步提高R-T-B系稀土類永久磁鐵之 等磁特性。 有鑒於此’本發明之目的係提供作爲具有優異 的稀土類系永久磁鐵原料的1^1_：6系合金。 又’本發明之目的係提供上述由R_T_B系合金 R-T-B系稀土類永久磁鐵用微細粉末及r_t_b系稀 久磁鐵。 解決問題之手段 本發明者詳細觀察包含成爲R-T-B系稀土類永夕 之D y的R - T - B系合金組織’檢視組織之狀態與磁特 的關係。然後’本發明者確認下列事實，即當含Dy T-B系合金具有由R2t14B相所成之主相與R經濃縮；^ 相’以及Dy經濃縮之富Dy相（Dy-rich phase)時， R-T-B系合金薄片製得之微細粉末成形燒結而獲得之 B系稀土類永久磁鐵會成爲保磁力等磁特性優異之牧 完成本發明。 亦即本發明係提供下列各發明。 (1)—種R-T-B系合金’其係稀土類系永久磁 之原料，係至少包含Dy之R-T-B系合金（但r爲S 、La、Ce、Pr、Nd、Pm、Sm、Eu、Gd、Tb、Ho、 Tm、Yb、Lu中的至少一種，T爲包含 80質量％以 的過渡金屬，B爲包含50質量％以上B’且另包含0 %以上未滿5 0質量％之C、N中的至少一種者），其 磁力 特性 作之 類永 .磁鐵 ^性間 之R-:富R 將該 R-T- I，而 :鐵用 ；c ' Y Er、 上Fe 質量 :特徵 200903532 爲：具有表現r2t14b相等磁性用之主相，及與合金整體 的組成比相較R係經濃縮之富R相，以及在前述富R相 附近形成，且與前述組成比相較Dy係經濃縮之Dy濃縮 區域。 (2 ) ( 1 )記載之R-T-B系合金，其中，Dy之濃度 係前述主相者低於前述Dy濃縮區域，前述富R相者低於 前述主相。 (3) 前述（1)或（2)記載之R-T-B系合金，其係 以薄帶連鑄法製得之平均厚度〇.1至1mm之薄片。 (4) 前述（1)至（3)中任一項記載之R-T-B系合 金之製造方法，其特徵在於隨著成爲平均厚度0.1至lmm 薄片的同時，往冷卻滾筒之平均熔液供給速度爲每lcm寬 每秒1 0 g以上。 (5 )前述（4 )記載之R-T-B系合金之製造方法，其 特徵在於將脫離冷卻滾筒的R-T-B系合金之薄片用6〇〇至 900 °C保溫30秒以上。 (6 )—種R-T-B系稀土類磁鐵用微細粉末，其係由 依據前述（1)至（3)中任一項記載之R-T-B系合金，或 (4 )或（5 )記載的R-T-B系合金之製造方法製得之r_ T-B系合金所製得者。 (7 ) —種R-T-B系稀土類永久磁鐵，其係由（6 )記 載之R-T-B系稀土類永久磁鐵用微細粉末製得者。 發明之效果 -9- 200903532 本發明之R-Τ-Β系合金，因具有在富R相附近形成且 與組成比相較Dy係經濃縮之Dy濃縮區域，故成爲可實 現保磁力高之磁特性優異的稀土類永久磁鐵。 再者，本發明之R-Τ-Β系稀土類永久磁鐵用微細粉末 及R-Τ-Β系稀土類永久磁鐵，由於是以本發明之R-Τ-Β系 合金或本發明之R-Τ-Β系合金之製造方法製得之R-Τ-Β系 合金所製得者，所以成爲保磁力高而磁特性優異者。 實施發明之最佳形態 第1圖表示本發明R-Τ-Β系合金之一例的照片，係以 掃描型電子顯微鏡（SEM)觀察R-Τ-Β系合金之薄片之剖 面時的照片。又，第1圖中左側爲鑄型面側。 第1圖所示之R-Τ-Β系合金係以SC法製得者。此R-T · B系合金之組成係質量比爲N d 2 3 %、D y 9 %、B 1 %、C 〇 1°/。' Ga 0.2%、其餘爲Fe。又，本發明之R-Τ-Β系合金之 組成並非限定於上述範圍，只要至少包含Dy之R-Τ-Β系 合金（但 R 爲 Sc、Y、La、Ce、Pr、Nd、Pm、Sm、Eu、 Gd、Tb、Ho、Er、Tm、Yb、Lu中的至少一種，T爲包含 8 0質量％以上F e的過渡金屬，B爲包含5 0質量％以上B ，且包含0質量％以上未滿5 0質量％之C、N中的至少一 種者），則任何組成均可。 第1圖所示之R-Τ-Β系合金係由r2t14B相（主相） 與富R相組成。第1圖中，富R相以白色表示，R2T14B 相（主相）以灰色表示。R2T14B相主要由柱狀晶、部分由 -10- 200903532 等軸晶所形成。R2 Τ !4 B相之短軸方向之平均結晶粒徑爲 10至50/zm。在R2T“B相之晶界與晶粒內，存在著延 R2TMB相之柱狀晶之長軸方向伸展之線狀富r相或—部 份中斷而成粒狀之富R相。富R相與組成比相較係r經 濃縮之非磁性且低熔點之相。富R相之平均間隔爲3至 1 0 // m。 第2圖至第6圖顯示藉由第1圖所示之R-T-B系合金 之電子探針微分析儀（ΕΡΜΑ)之波長分散型X射線分光 儀（WDS )的元素分佈分析（數字側繪圖；Digital mapping ) 之結果。 第2圖爲第1圖所示之R-T-B系合金之電子束圖像， 富R相以白色表示，R2T14B相（主相）以灰色表示。 第3圖表示對應第2圖區域之Fe之分佈圖。從第2 圖及第3圖中可知富R相與主相相較，富R相之F e較少 〇 第4圖表示對應第2圖區域之Nd2分佈圖。從第2 圖及第4圖中可知富R相與主相相較’富R相之Nd較多 〇 第5圖表示對應第2圖區域之Dy之分佈圖。從第2 圖及第5圖中可知富R相與主相相較’富R相之Dy較少 〇 第6圖表示對應第2圖區域之〇3之分佈圖。從第2 圖及第6圖中可知富R相與主相相較’富R相之Ga較多 -11 - 200903532 又，第7圖至第10圖顯示藉由場發射型電子微探儀 (FE-EPMA)之元素分佈分析（數値側繪圖’ DigitaI m a p p i n g )之結果者。 第7圖爲第1圖所示之R-T-B系合金之電子束圖像’ 富R相以白色表示，R2T14B相（主相）以灰色表示。 第8圖表示對應第7圖區域之Dy之分佈圖。從第7 圖及第8圖可知富R相與主相相較，Dy經濃縮之Dy濃縮 區域係在富R相附近形成。再者，由第8圖可知Dy之濃 度係主相少於Dy濃縮區域，而富R相又少於主相。 又，第9圖表示對應第7圖區域之Fe之分佈圖。從 第7圖及第9圖中可知富R相與主相相較，富R相之Fe 較少。 第10圖表不對應第7圖區域之Nd之分佈圖。從第7 圖及第10圖中可知富R相與主相相較，富R相之！^(1較 多。 (製造方法） 第1圖所示本發明之R-T-B系合金，舉例如，可藉由 使用第11圖所示之合金製造裝置之SC法鑄造。 「合金製造裝置」 第1〗圖表示本實施形態之合金製造裝置的整體構造 之正面模式圖。 第11圖所示之合金製造裝置丨（下文以製造裝置】表 -12- 200903532 示），大致由鑄造裝置2、破碎裝置21、加熱裝置3所構 成。加熱裝置3大致由加熱器31及容器5所構成。容器5 大致由儲藏容器4、以及設置於儲藏容器4上方之開關式 平台（stage )組32所構成。 第11圖所不之製造裝置1備置箱室（chamber) 6。 箱室6係由鑄造室6a、以及設置於鑄造室6a下方且與鑄 造室6a連通之保溫·儲藏室6b所構成。在鑄造室6a收 纳鑄造裝置2’在保溫·儲藏室6b內收納加熱裝置3。又 ，在保溫·儲藏室6b設置閘門（gate ) 6e，除了將容器5 運送至保溫·儲藏室6b外部外’保溫•儲藏室6b藉由閘 門6 e而密閉。 另，於鑄造裝置2備置破碎裝置21，在鑄造裝置2及 開關式平台組3 2之間備置將鑄造合金薄片引導至開關式 平台組32之斗槽7。 「鑄造裝置」 第12圖爲裝備於製造裝置1內之鑄造裝置2之正面 模式圖。 第1 2圖所示之鑄造裝置2 ’係大致由以未圖示的水冷 機構急速冷卻合金熔液L而鑄造鑄造合金μ的冷卻滾筒 2 2 '及將合金熔液L供給至冷卻滾筒2 2之澆鑄分配器（ tun di sh ) 23、以及將鑄造合金μ破碎而做成鑄造合金薄 片Ν之破碎裝置21所構成。破碎裝置21如第12圖所示 ’例如由一對破碎滾筒2 1 a所構成。 -13- 200903532 「加熱裝置」 第13圖表示裝備於製造裝置1內之加熱裝置3之正 面模式圖，第14圖爲側面模式圖’第15圖爲平面模式圖 〇 如第13圖至第15圖所示’構成加熱裝置3之加熱器 31，係由加熱器蓋31a、及安裝於加熱器蓋31a下方之加 熱器本體31b所構成。加熱器蓋31a係用於使加熱器本體 31b產生的熱放射至容器5側面’且防止加熱器本體31b 產生的熱放射至鑄造室6a而設計。又’藉設置加熱器蓋 3 1 a，而可防止合金熔液或鑄造合金之一部分從鑄造裝置2 落下時加熱器本體31b之破損。 又，在加熱器3 1處設置開口部3 1 c，在此開口部3 1 c 配置斗槽7之排出口 7a。藉此，可將通過斗槽7而從鑄造 裝置2落下來之鑄造合金薄片N供給至加熱器31下方之 容器5的開關式平台組3 2。 再者，如第1 1圖及第13圖所示，沿著於保溫·儲藏 室6b內設置之傳送帶51的較長方向（容器5之移動方向 )配置加熱器3 1。 藉由此構成，即使於保溫•儲藏室6b內部移動容器5 時，裝載於容器5開關式平台組32之鑄造合金薄片N仍 係均勻的保溫。 其次’構成加熱裝置3之開關式平台組3 2，與儲藏容 器4形成一體而構成容器5。即第13圖至第15圖所示之 -14- 200903532 容器5係由儲藏容器4、及設置於儲藏容器4上方 式平台組32所構成。 於開關式平台組3 2 ’沿谷:5的移動方向序列 複數開關式平台33。且，在開關式平台組32周圍 向構件52’藉由此導向構件52防止通過斗槽7落 造合金薄片N散亂於保溫·儲藏室6b內。 各開關式平台3 3係將禱造裝置2供給·之|壽造 片N載置至以加熱器31預定之保溫時間。在歷經 間後使鑄造合金薄片N落至儲藏容器4者。在各開 台33，分別備置各種平台板33a、及開關平台板3: 關結構3 3 b。各開關結構3 3 b係分別由安裝於平台 —側之迴轉軸3 3 b】、及旋轉驅動此迴轉軸3 3 b 1未 各驅動源所構成。各平台板3 3 a之傾斜角度，藉各 可各別控制迴轉軸3 3 b i之旋轉。各平台板3 3 a之 度得以在〇 ° (平台板3 3 a爲水平狀態（第1 3圖中 虛線表示之狀態））至順時針方向約9 0° (平台板 乎爲垂直狀態（第13圖中以實線表示之狀態）） 間任意設定。 又，開關式平台33係藉操作開關結構33b，在 預定保溫時間之前將鑄造合金薄片N載置於平台板 之後使平台板3 3 a之傾斜角度變大而使鑄造合金薄 入儲藏容器4內。 另外，開關式平台3 3係藉由發揮作爲儲藏容； 蓋的任務，使加熱器3 1之熱度不會到達儲藏容器4 之開關 備置之 設置導 下的鑄 合金薄 保溫時 關式平 3 a之開 板33a 圖示之 驅動源 傾斜角 以兩點 33a幾 之範圍 至經過 :33a， 片N落 器4之 μ可防 -15- 200903532 止儲藏容器4內部之溫度上昇。又，儲藏容器4之內部係 放置多片冷卻板4a。 又’如第1 3圖及第14圖所示’容器5係裝載於傳送 帶51 (可動裝置）上方。容器5藉由傳送帶51而可在第 13圖中往左右方向移動。 「合金之鑄造」 第16圖至第19圖之任一個皆爲說明合金製造裝置之 操作之正面模式圖。 首先’如第1 6圖所示’使位於開關式平台組3 2中的 圖中左端之開關式平台33A處於斗槽7之排出口 7a之正 下方而使容器5移動。且’將全部的開關示平台33調至 「關閉」的狀態。 其次，使第12圖之鑄造裝置運作而製作鑄造合金薄 片N。首先，在未圖示之熔解裝置中調製合金熔液L。合 金熔液L之溫度係根據合金成分而調整至130(TC至1500 °C之範圍。製得之合金熔液L係連同耐火物坩堝24搬送 至鑄造裝置2 ’從耐火物坩堝24供給至澆鑄分配器23。 然後從澆鑄分配器2 3供給至冷卻滾筒2 2，在冷卻滾筒2 2 上凝固而做成鑄造合金Μ。之後，使鑄造合金Μ在澆鑄分 配器23之反面側自冷卻滾筒22脫離’將鑄造合金Μ挾持 於2個旋轉之破碎滾筒2 1 a間使之破碎’製作鑄造合金Ν 〇 供至冷卻滾筒2 2之平均熔液供給速度係每1公分寬 -16- 200903532 可達到每秒1 〇g以上，以每1公分寬可達到每秒20g以上 爲佳，以每1公分寬可達到每秒25g以上更佳，每1公分 寬可達到每秒1 00g以下者又更佳。合金熔液L之供給速 度低於每秒1 〇g時，由於合金熔液L自身之黏性或其與冷 卻滾筒22表面之潤濕性，合金熔液L無法薄薄地廣泛潤 濕於冷卻滾筒22上而會收縮，造成合金品質之變動。又 ，供至冷卻滾筒22之平均熔液供給速度每1公分寬超過 每秒1 00g時，冷卻滾筒22上之冷卻會不充足而產生組織 粗大化、a -Fe析出之現象。 又，冷卻滾筒22上之合金熔液的平均冷卻速度以每 秒100至2 000 °C爲佳。若每秒100°C以上則冷卻速度充足 ，且可防止a -Fe之析出、富R相等組織之粗大化。另外 ，若每秒2 00(TC以下，則過冷度不會變得過剩，而可以適 當溫度將鑄造合金薄片供給至加熱裝置3。再者，由於鑄 造合金薄片不會過冷，所以沒有必要再加熱。平均冷卻速 度可由熔液剛接觸冷卻滾筒時之溫度與脫離冷卻滾筒時之 溫度差除以接觸冷卻滾筒上之時間而求得。 再者，脫離冷卻滾筒22時，鑄造合金Μ之平均溫度 會殷鑄造合金Μ與冷卻滾筒22接觸程度之微妙差異、鑄 造合金Μ厚度之不勻等而微妙地變化。鑄造合金Μ脫離 冷卻滾筒之平均溫度，例如可利用放射溫度計於鑄造開始 至終止時，於寬度方向掃描合金表面而測定，並將得到之 測量値平均化而求得。 鑄造合金Μ脫離冷卻滾筒22之平均溫度以較合金熔 -17- 200903532 液之r2ti4b相平衡狀態下之凝固溫度低100至500°c爲佳 ，低100至400t爲更佳。R2TI4B相之熔解溫度’ Nd_Fe· B三元系是設定在1 1 5 0 °C，但依據N d以之外稀土類元素 之置換、Fe以外之過渡元素之置換、其他添加元素之種類 或添加量而變化。脫離冷卻滾筒22之鑄造合金M之平均 溫度與鑄造合金Μ的R2T, 4B相平衡狀態下的凝固溫度間 之溫差，若未滿loot：時’等同冷卻不足。另一方面’其 溫差超過5 0 0 °c時，因冷卻速度過快而形成熔液之過冷。 另外，脫離冷卻滾筒22之鑄造合金Μ之平均溫度’ 即使在同一鑄造步驟（出渣（tap ))內也會變動’但若 其變化幅度大則會造成組織、品質之變動。因此’出鐵內 之溫度變化幅度，以小於200 °C爲宜，較佳者爲1〇〇 °C以 下，更佳者爲50°C，又更佳者爲20°C。 鑄造合金薄片N之平均厚度以成爲0.1mm以上1mm 以下爲佳。薄片之平均厚度若比0.1mm薄則凝固速度過度 增加，富R相之分散會變得過細。又，薄片之平均厚度若 比1 mm厚則會因凝固速度降低導致富R相之分散性降低 、a - F e之析出等。 其次，如第16圖所示，鑄造合金薄片N通過斗槽7 內送至加熱裝置3，堆積（載置）於位於斗槽7之排出口 7a正下方的開關式平台33 A上。此時，加熱器3 1處於通 電狀態’鑄造合金薄片N載置於開關式平台3 3 A上之後 ’隨即藉由加熱器3 1保溫或升溫。 對開關式平台33A而言，鑄造合金薄片n之堆積量 -18- 200903532 只要依據平台板3 3 a之面積適當設定即佳，但因鑄造 薄片N係從鑄造裝置2連續供應，雖根據供給速度等 任一情況其鑄造合金薄片N皆會由開關式平台3 3 A 。因此，對應於開關式平台33A之鑄造合金薄片N 積量達設定値時’如第17圖所示將容器5往圖中左 移動’使位於開關式平台3 3 A右邊之開關式平台3 3 B 斗槽7排出口 7a之正下方，使鑄造合金薄片n堆積 開關式平台33B。之後，同樣操作’—邊配合鑄造合 片N之調製而移動容器5，一邊依次使鑄造合金薄片 積於各開關式平台33C至33E。 堆積於各開關式平台33A至33E之鑄造合金薄 分別藉由加熱器3 1保溫或升溫。保溫溫度以較滾筒 溫度低爲佳’具體而言以（滾筒脫離溫度—1 〇 〇。<：) 至滾筒脫離溫度以下之範圍爲佳，（滾筒脫離溫度-)以上滾筒脫離溫度以下之範圍爲佳，更具體而言以 °C以上900°C以下之範圍爲佳。若保溫溫度爲6〇〇。〇 ’則可充分提高R-T-B系合金之保磁力。又，若保溫 爲9 0 0 °C以下’則可防止a - F e之析出、富R相等組 粗大化。 另外，若因任何理由而滾筒脫離溫度降低時，亦 保溫溫度昇溫至較滾筒脫離溫度更高而保持之。較佳 幅度以100°c以內爲佳’ 50°c以內更佳。昇溫幅度若 時生產效率降低。又’即使在1 0〇ot保溫，亦有提昇 力之效果。但，組織粗大化，微細粉碎時之粒度分佈 合金 ，然 溢出 之堆 方向 移至 於此 金薄 N堆 片N 脫離 以上 5 0°C 600 以上 溫度 織之 可將 昇溫 過闻 保磁 及微 -19- 200903532 細粉末之流動性，甚至燒結溫度都會變化。因此，在1 〇 〇 〇 °c保溫時，必須考慮對後續步驟之影響。 再者，保溫時間以3 0秒以上爲佳，3 0秒至數小時左 右更佳，3 0秒至3 0分鐘左右最佳。保溫時間若爲3 0秒以 上者即可充分提高保持力，亦可保溫數小時，但就生產效 率方面而言以30分鐘以下爲宜。 其次，如第18圖所示，關於其餘之開關式平台33F 至33J，同樣地亦可藉由配合鑄造合金薄片N之調製使容 器5移動，而依次使鑄造合金薄片N堆積於各開關式平台 33F至33J。又堆積於開關式平台33A至33D之鑄造合金 薄片N各自若經過預定之保溫時間，則如第1 8圖所示， 使各開關式平台33A…依次成爲「開」之狀態，使鑄造合 金薄片N依次落至儲藏容器4。藉由使鑄造合金薄片N落 下儲藏容器4，加熱器31之熱度不會到達鑄造合金薄片N ，藉此終止保溫處理。 如第1 7圖中所說明，由於將鑄造合金薄片N依次載 置於各開關式平台33A…上，所以對各開關式平台33A… 上之鑄造合金薄片N之保溫起始時間在各開關式平台3 3 每個都有時間差。因此，爲了使對各開關式平台33A…上 之鑄造合金薄片N之保溫時間維持一定，以使各開關式 33A…依次成爲「開」之狀態，使鑄造合金薄片N依次落 至儲藏容器4爲佳。 落於儲藏容器4之鑄造合金薄片N因接觸冷卻板4a 而熱度被冷卻板4a吸取，藉此冷卻鑄造合金薄片N。 -20- 200903532 第19圖及第20圖顯示全部之開關式平台33A…成爲 「開」之狀態，而鑄造合金薄片N收納於儲藏容器4之狀 態。其後，若繼續以鑄造裝置2進行鑄造、破碎步驟時， 則使全部開關式33A…成爲「閉」之狀態，並一邊使容器 5往圖中右方向移動，一邊配合鑄造合金薄片N之調製使 鑄造合金薄片N依次載置於各開關式平台33 A…上即可。 又，使鑄造裝置2之鑄造、破碎步驟終止時，則使全部開 關式33 A…成爲「閉」之狀態而使加熱器31之熱度不會 到達儲藏容器4。然後，打開保溫·儲藏室6b之閘門6e 並將容器5運送至室6外部，取出鑄造合金薄片N，終止 鑄造合金薄片N之製造。 「冷卻速度」 其次說明關於製造R-T-B系合金時之冷卻速度。 本發明中，於剛凝固之溫度之主相凝固點（1 1 70 °C附 近）至較富R相凝固點的約700 °C更低的600 °C爲止，將 其冷卻速度控制使成爲如下所示之冷卻速度。 R_T-B系合金之冷卻速度在1 000 °C至8 5 0 °C間爲每秒 100至3 00 °C。l〇〇〇°C至8 5 0 °C之冷卻速度，若較上述範圍 更快時則慮及Dy不會充分擴散至主相，若較上述範圍更 慢時則慮及Dy會過剩擴散而無法形成主相中之濃縮部分 〇 又，R-T-B系合金之冷卻速度，從主相凝固點至1〇〇〇 °C之間以成爲每秒3 00至2000 °C爲宜。由於使主相凝固點 -21 - 200903532 至1000 °c之冷卻速度成爲上述範圍，可獲得具有Dy濃縮 區域之R-T-B系合金，且可獲得高生產力。 又，R-T-B系合金之冷卻速度，在850。(：至600°C區 域暫時要求以成爲每秒1 0 0 °C以下爲宜。由於暫時使8 5 0 C至600C之冷卻速度成爲上述範圍，可充分使富r相所 含之Dy擴散至鄰接之主相，所以可容易地製造具有Dy ί辰縮區域而保fe力更商之R-T-B系合金。 由於本實施形態之R-T-B系合金及R_T_B系合金薄片 係於富R相附近形成’且具有與組成比相較Dy經濃縮之 Dy濃縮區域，所以成爲可實現具有高保磁力、而磁特性 優異之稀土類永久磁鐵者。 亦即’本實施形態之R-T-B系合金，例如具有較第 21圖至第24圖所示不具有Dy濃縮區域之R-T-B系合金 更高的保磁力。 第21圖至第24圖顯示以不具有Dy濃縮區域之R-T-B系合金之一例的FE-EPMA (場發射型電子微探儀）元素 分佈分析之結果。第21圖至第24圖所示之R_T_B系合金 係用SC法所製得者。此R-T-B系合金之組成爲質量比： N d 2 3 %、D y 9 %、B 1 %、C ο 1 %、g a 0 · 2 %，其餘部分爲 F e ° 第21圖爲不具有Dy濃縮區域之r_T_b系合金之電子 束圖像’富R相以白色表示，r2t14b相（主相）以灰色 表示。 第22圖表示對應第21圖區域之Dy之分佈圖。從第 -22- 200903532 2 1圖及第2 2圖可知與主相相較並無D y經濃縮之D y濃縮 區域，而Dy之濃度與主相相較，富R相亦少。 第23圖表示對應第21圖區域之Fe之分佈圖。從第 21圖及第23圖中可知與主相相較’富R相之Fe較少 第24圖表示對應第21圖區域之Nd之分佈圖。從第 21圖及第24圖可知與主相相較，富R相之Nd較多。 (R-T-B系稀土類永久磁鐵之製作） 在製作本發明之R - T - B系稀土類永久磁鐵方面’首先 由本發明之R-T-B系合金製作R-T-B系稀土類永久磁鐵用 微細粉末。本發明之R-T-B系稀土類永久磁鐵用微細粉末 係例如藉由使本發明之R-T-B系合金所成之薄片吸收氫’ 經氫解碎後使用噴射硏磨機等粉碎機進行微細粉碎之方法 而獲得。此處的氫解碎係例如以預先進行保持於預定壓力 之氫氣氛圍中的氫吸收步驟爲宜。 其次，將所獲得之R-T-B系稀土類永久磁鐵用微細粉 末，例如使用橫磁場中成型機等進行壓製成形並於真空中 燒結而獲得R-T-B系稀土類永久磁鐵。 本實施形態之R-T_B系稀土類永久磁鐵用微細粉末及 R-T-B系稀土類永久磁鐵由於是由本發明之R-T_B系合金 製得者，因而係保磁力高、磁特性優異者。 【實施方式】 〔實施例1〕 -23- 200903532 將依質量比爲：N d 2 3 %、D y 9 %、B 0 · 9 8 %、C ο 1 %、 G a 0 · 2 %、及其餘部分爲F e加以配合之原料稱量，並使用 氧化鋁坩堝，在氬氣1氣壓氛圍中以高周波熔解爐熔解而 調製合金熔液。接著’將此合金熔液供給至第1 1圖所示 製造裝置的鑄造裝置’用SC法鑄造。鑄造時冷卻滾筒之 周速度爲1 .3m/s，供給至冷卻滾筒之平均熔液供給速度爲 每1公分寬每秒3 0 g ’鑄造合金塊脫離冷卻滾筒之平均溫 度爲8 5 0 C。 此合金之冷卻速度爲：從主相之凝固點至1 〇 〇 〇 °c爲 700°C /秒，1 000°C 至 850°C 爲 200°c /秒，8 5 0 °C 至 7 80 °c 爲 5 (TC /秒，其後，使用第11圖之製造裝置於開關平台上以 78 0 °C左右之溫度保持3 00秒，之後以0.1°C /秒之冷卻速 度冷卻至6〇〇°C以下，製作實施例1之r-Τ-Β系合金薄片 。此時合金之平均厚度爲〇.3mm。 〔實施例2〕 使用與實施例1相同之原料與裝置製作合金熔液。然 後，使用與實施例1相同之鑄造裝置，將所獲得之合金熔 液於設定爲：鑄造時之冷卻滾筒周速度〇 _ 8 7 m /秒、供給至 冷卻滾筒之平均熔液供給速度每〗公分寬每秒3〇g、鑄造 合金塊脫離冷卻滾筒之平均溫度爲8 8 〇°C之條件下進行鑄 造。 此合金之冷卻速度爲：從主相之凝固點至1 0 0 〇 °c爲 7 00 °C /秒，1 000 °C 至 850°c 爲 2 0 0 °C /秒，8 5 0°C 至 7 80°c 爲 -24- 200903532 10C/秒，其後’使用桌11圖之製造裝置，不使用開關平 台以0_l°c/秒之冷卻速度冷卻至600 °c以下，製作實施例 2之R-T-B系合金薄片。此時之合金平均厚度爲〇45mm。 關於所獲得之實施例1及實施例2之R_T_B系合金薄 片，以WDS (波長分散形X射線分光儀）.EpMA及FE-EPMA之元素分佈分析（數字側繪圖；Digital mapping ) (表面分析）。結果，實施例1及實施例2之R-T-B系合 金薄片’任一者均在富R相附近’形成富R相及與主相相 較’ D y經濃縮之D y濃縮領域。另外，實施例1及實施例 2之R-T-B系合金薄片，任一者之Dy濃度均爲主相少於 Dy濃縮區域，而富R相更少於主相。 (比較例1 ) 使用與實施例1相同之原料與裝置製作合金熔液。使 用與實施例1相同之鑄造裝置，將所獲得之合金熔液於設 定爲：鑄造時冷卻滾筒之周速度0.65m/s、冷卻滚筒之平 均熔液供給速度每1公分寬1 5 g/秒、而鑄造合金塊脫離冷 卻滾筒之平均溫度爲700°C之條件下進行鑄造，製得比較 例1之R-T-B系合金薄片。 此合金之冷卻速度爲：從主相之凝固點至1 〇〇〇 °C爲 700°C/秒，l〇〇〇°C 至 700°C 爲 400°C/秒，7〇〇°C 至 600 °C 爲 j 〇r /秒，其後，使用第11圖之製造裝置而不使用開關平 台，以每秒0.1 °C之冷卻速度冷卻至600 °C以下。此時合 金之平均厚度爲〇.30mm。 -25- 200903532 關於所獲得之比較例1之R-Τ-Β系合金薄片 WDS_EPMA及FE-EPMA進行元素分佈分析（數字側 ；Digital mapping )(表面分析）。結果，比較例1 [B系合金薄片，並未形成與組成比相較，Dy經濃 D y濃縮區域。咸認爲其中一個原因係於比較例1中 於鑄造合金塊脫離冷卻滾筒之溫度低，合金在冷卻滾 過度急速冷卻，且I OOOt至70 0 t之冷卻速度過快， Dy及Nd無法充分擴散而無法形成濃度梯度。 接著，如下所示，使用所獲得之實施例1及實施 、比較例1之R-Τ-Β系合金薄片製作磁鐵。 首先，將實施例1及實施例2、比較例1之R-T 合金薄片進行氫解碎。氫解碎係藉由使各R-Τ-Β系合 片在2氣壓之氫中吸收氫後，在真空中加熱至500 °C 殘留的氫，之後添加0.07質量％硬脂酸鋅並使用氮氣 射硏磨機加以微細粉碎之方法進行。經微細粉碎而獲 粉未以雷射繞射式測量之平均粒度爲約5.0 m。 其次，使用橫磁場中成型機將獲得之粉末體在 氮氣氛圍中以〇.8t/cm2成形壓力壓製成型而獲得成形 而後，在1.33xl0_5hPa之真空中將獲得之成形體從室 昇溫，於500°C、8 00°C各保持一小時，以去除硬脂酸 殘留之氫。其後，升溫至燒黏溫度之1 〇 3 0 °C，保持3 而製得燒黏體。之後，將獲得之燒黏體於氬氣氛圍中 8 0 0 °C、5 3 0 °C各別進行熱處理一小時，於實施例1至 例2中各獲得I 〇個，比較例獲得5個磁鐵。 ，以 繪圖 之R- 縮之 ，由 筒上 導致 例2 -B系 金薄 抽出 流噴 得之 1 0 0 % 體。 溫起 鋅及 小時 ，於 實施 -26- 200903532 而後’以B Η磁滯曲線儀（B H C u r v e T r a c e r )測量所 獲得之實施例1及實施例2，比較例1之磁鐵之磁特性。 結果如表1及第25圖所示。第25圖爲顯示實施例1、實 施例2及比較例1之磁鐵之保磁力（η Cj )的圖表，縱軸 表示保磁力’橫軸表式水準。又，第25圖中符號〇表示 實施例1及實施例2之保磁力，符號▲表示比較例1之保 磁力。 -27- 200903532 〔表1〕200903532 IX. INSTRUCTIONS OF THE INVENTION [Technical Field] The present invention relates to a fine powder of an R-Τ-lanthanide alloy 'R-Τ-lanthanide rare earth permanent magnet, and an RTB rare earth permanent magnet, in particular, A fine powder of R-Τ-lanthanide alloy and R-Τ-lanthanide rare earth permanent magnet of R-Τ-lanthanide rare earth permanent magnet excellent in magnetic force. [Prior Art] R-Τ-Β-based magnets are often used in hard disk (HD), magnetic resonance imaging (MRI), various motors, etc. due to their high characteristics. In recent years, as the heat resistance of the R_T-B magnet has increased and the demand for energy conservation has increased, the motor use ratio including the automobile has increased. R-Τ-Β-based magnets are collectively referred to as Nd-Fe-B or R-Τ-Β-based magnets because their main components are Nd, Fe, and B. A part of the R-based Nd of the R-Τ-lanthanum magnet is replaced with another rare earth element such as Pr, Dy or Tb. τ is one of Fe and is replaced by other transition metals such as Co and Ni. B is boron, and a part thereof may be replaced with C or N. The R-T-B alloy of the R-T-B magnet is a main phase of the R2TMB phase which contributes to the magnetization of the magnetic phase, and the alloy of the non-magnetic rare earth element and the low melting point of the R phase. . Since R-Τ-lanthanum alloys are generally activated or melted in a vacuum or an inert gas because of their active metals. Further, in the case of producing a sintered magnet by a powder metallurgy method for a cast R-τ B-based alloy, the alloy block is generally pulverized to an average particle diameter of about 5 μm (d50: measured by a laser diffraction type particle size distribution meter) to form an alloy powder. After that, in the magnetic-4 - 200903532 field, extrusion molding is performed in a sintering furnace at a high temperature of about 10,000 to 1,100 ° C, and then heat treatment and mechanical processing as needed to further improve corrosion resistance. Plated to form a sintered magnet. In the R-T-B based sintered magnet, the R-rich phase plays an important role as described below. 1) The melting point is low, the liquid phase is formed during sintering, and the magnet is increased in density, thereby contributing to the improvement of magnetic properties. 2) The unevenness of the grain boundary is eliminated, and the nucleation point (nucleation s i t e ) of the reverse magnetic field is reduced to increase the coercive force. 3) Magnetically insulate the main phase to increase the coercive force. Therefore, if the dispersion state of the R-rich phase in the formed magnet is poor, local sintering failure and magnetic deterioration are caused, and it is therefore important that the R-rich phase is uniformly dispersed in the formed magnet. The distribution of the R-rich phase of the R-T-B sintered magnet is greatly affected by the structure of the raw material R-T-B alloy. Further, another problem derived from the casting of R-T-B alloys is the formation of a-Fe in the cast alloy. Since α-Fe has a deforming force and remains in the pulverizer without being pulverized, not only the pulverization efficiency at the time of pulverizing the alloy is lowered, but also the composition variation and the particle size distribution before and after the pulverization are affected. Further, if the magnet remains in the magnet after the a-Fe is sintered, the magnetic properties of the magnet are weakened. Therefore, the conventional alloys are required to be subjected to a long-term homogenization treatment at a high temperature to eliminate a-F e . However, since α _ ρ e exists as a peritectic nucleus, its elimination must be a solid phase diffusion for a long time. If it is an ingot of several cm thickness and the rare earth amount is 33% or less, it is virtually impossible to eliminate it. α -Fe. 200903532 In order to solve the problem of α-Fe production in this R-Τ-lanthanum alloy, a strip casting method (referred to as SC method) for casting an alloy block at a rapid cooling rate has been developed and put into practical use. The SC method is a method in which the alloy is rapidly solidified by a sheet which is poured into a molten metal on a water-cooled copper cylinder to a thickness of about 1 mm. In the case of the SC method, since the molten metal is excessively cooled to a temperature lower than the formation temperature of the main phase R2T14B phase, the R2T14B phase can be directly formed from the alloy melt, and the deposition of a-Fe can be suppressed. Further, by performing the S C method, since the crystal structure of the alloy is made fine, an alloy having a structure in which the R-rich phase is finely dispersed can be produced. The R-rich phase reacts with hydrogen in a hydrogen atmosphere to become an expanded and brittle hydride. If this property is utilized, fine cracks corresponding to the degree of dispersion of the R-rich phase can be introduced. When the fine pulverization is carried out after the hydrogenation step, the pulverizability is excellent because the alloy is destroyed by a large amount of fine cracks formed by hydrogenation. Thus, the alloy cast by the SC method has a fine dispersion of the R-rich phase, so that the dispersibility of the R-rich phase in the pulverized and sintered magnet is also good, and the magnetic properties of the magnet are successfully improved (for example, refer to the patent). Document 1). Further, the alloy flakes cast by the SC method have good homogeneity in structure. The homogeneity of the structure can be compared by the crystal grain size or the dispersion state of the R-rich phase. The alloy flakes obtained by the SC method produce chi-cooled crystals (chi 11 crysta 1 ) on the side of the alloy flake casting drum (hereinafter also referred to as the mold surface side), but can be moderately caused by the overall rapid cooling solidification. Fine and homogeneous organization. As described above, the R-T-B alloy cast by the SC method has an excellent structure for producing a sintered -6 - 200903532 magnet because the R-rich phase is finely dispersed and the formation of a-Fe is also suppressed. The magnetic properties, particularly the relationship between the coercive force and the elemental distribution in the fine structure of the magnet, have a great influence on the distribution of D y which contributes to the improvement of the coercive force. For example, it has been reported that if Dy is distributed in the vicinity of the grain boundary phase, the coercive force is increased (for example, refer to Patent Document 2). More specifically, it has been reported that the coercive force is increased when these Dy are present in the main phase (see, for example, Patent Document 3 and Non-Patent Document 1). Moreover, since the characteristics of the magnet have a certain correlation with the alloy manufacturing method, the manufacturing method of the alloy is also improved as the magnet characteristics are improved. For example, a known method of controlling the fine structure (for example, refer to Patent Document 4), and a method of controlling the surface structure of the casting drum to a predetermined thickness to control the fine structure (for example, refer to Patent Document 5 and Patent Document 6). [Patent Document 1] Japanese Laid-Open Patent Publication No. Hei No. 5-2 2 2 8 8 8 (Patent Document 2) Japanese Patent Laid-Open No. Hei 5 _ 2 1 2 1 9 Patent Publication [Patent Document 3] W02003/001541 [Patent Document 4] [Patent Document 5] Japanese Patent Laid-Open No. 2 〇 0 3 _ 丨 8 8 8 8 专利 专利 专利 专利 专利 专利 专利 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富 富Body and Powder Metallurgy, Vol. 52, No. 3, March 2005, pp. 158-163. SUMMARY OF THE INVENTION Problems to be Solved by the Invention However, in recent years, R_T_B rare earth permanent magnets for higher performance have been sought. 200903532 Iron, it is necessary to further improve the magnetic properties of R-T-B rare earth permanent magnets. In view of the above, an object of the present invention is to provide a 1^1_:6 alloy which is an excellent rare earth permanent magnet raw material. Further, the object of the present invention is to provide the fine powder for R-T_B-based alloy R-T-B rare earth permanent magnet and the r_t_b-based rare magnet. MEANS FOR SOLVING THE PROBLEMS The inventors of the present invention have observed in detail the relationship between the state of the structure and the magnetic properties of the R-T-B alloy structure which is a D-y of the R-T-B rare earth. Then, the inventors confirmed the fact that when the Dy TB-based alloy has a main phase formed by the R2t14B phase and R is concentrated; ^ phase' and Dy are concentrated in a Dy-rich phase, RTB The B-based rare earth permanent magnet obtained by forming and sintering the fine powder obtained by the alloy sheet is excellent in magnetic properties such as coercive force and the like. That is, the present invention provides the following inventions. (1) An RTB-based alloy, which is a rare earth-based permanent magnetic material, which contains at least an RT-based alloy of Dy (but r is S, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb) And at least one of Ho, Tm, Yb, and Lu, T is a transition metal containing 80% by mass, B is 50% by mass or more, and further contains 0% or more and less than 50% by mass of C, N. At least one of them), its magnetic properties are like eternal. R-: between the magnets R: R rich for RT-I, and: for iron; c 'Y Er, upper Fe quality: characteristic 200903532 for: performance The main phase of r2t14b is equal to the magnetic phase, and the composition ratio of the whole alloy is compared with the R-rich phase of the R-rich phase, and is formed in the vicinity of the R-rich phase, and the Dy-concentrated Dy concentrated region is compared with the aforementioned composition ratio. . (2) The R-T-B alloy according to (1), wherein the concentration of Dy is such that the main phase is lower than the Dy concentration region, and the R-rich phase is lower than the main phase. (3) The R-T-B alloy described in the above (1) or (2), which is a sheet having an average thickness of 11 to 1 mm obtained by a strip casting method. (4) The method for producing an RTB-based alloy according to any one of the above (1) to (3), wherein the average melt supply speed to the cooling drum is the same as the sheet having an average thickness of 0.1 to 1 mm. Lcm width is more than 10 g per second. (5) The method for producing an R-T-B alloy according to the above (4), characterized in that the sheet of the R-T-B alloy which has been removed from the cooling drum is kept at 6 Torr to 900 ° C for 30 seconds or longer. (6) The RTB-based alloy according to any one of the above (1) to (3), or the RTB-based alloy according to (4) or (5), which is a fine powder for a rare earth magnet of the present invention. Produced by the r_ TB alloy obtained by the manufacturing method. (7) A rare earth permanent magnet of R-T-B type, which is obtained by using fine powder of R-T-B rare earth permanent magnet recorded in (6). EFFECTS OF THE INVENTION-9-200903532 The R-Τ-lanthanum alloy of the present invention has a magnetic field with high coercive force because it has a Dy concentration region which is formed in the vicinity of the R-rich phase and which is concentrated in comparison with the composition ratio Dy. A rare earth permanent magnet with excellent properties. Further, the fine powder for R-Τ-lanthanide rare earth permanent magnet of the present invention and the R-Τ-lanthanum-based rare earth permanent magnet are the R-Τ-lanthanide alloy of the present invention or the R- of the present invention. Since the R-Τ-lanthanum alloy obtained by the method for producing a ruthenium-bismuth alloy is obtained, it has a high coercive force and excellent magnetic properties. BEST MODE FOR CARRYING OUT THE INVENTION Fig. 1 is a photograph showing an example of the R-Τ-lanthanum alloy of the present invention, which is a photograph of a cross section of a sheet of an R-yttrium-lanthanum alloy observed by a scanning electron microscope (SEM). Further, the left side in Fig. 1 is the mold surface side. The R-Τ-lanthanum alloy shown in Fig. 1 is obtained by the SC method. The mass ratio of the composition of the R-T · B alloy is N d 2 3 %, D y 9 %, B 1 %, C 〇 1 °/. 'Ga 0.2%, the rest is Fe. Further, the composition of the R-rhenium-yttrium-based alloy of the present invention is not limited to the above range, and it is only required to contain at least R-Τ-lanthanide alloy of Dy (but R is Sc, Y, La, Ce, Pr, Nd, Pm, At least one of Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb, and Lu, T is a transition metal containing 80% by mass or more of F e , and B is 50% by mass or more and contains 0 mass. Any of the components C and N which are less than 50% by mass, and any composition may be used. The R-Τ-lanthanum alloy shown in Fig. 1 is composed of an r2t14B phase (main phase) and an R-rich phase. In Fig. 1, the R-rich phase is shown in white, and the R2T14B phase (main phase) is shown in gray. The R2T14B phase is mainly composed of columnar crystals and partially composed of -10 200903532 equiaxed crystals. R2 Τ !4 The average crystal grain size of the B-phase in the short-axis direction is 10 to 50/zm. In the grain boundary and grain of the B phase of R2T, there is a linear r-rich phase extending in the long axis direction of the columnar crystal extending in the R2TMB phase or a R-rich phase partially broken into a granular shape. Compared with the composition ratio, the non-magnetic and low-melting phase is concentrated. The average interval of the R-rich phase is 3 to 10 // m. Figures 2 to 6 show the RTB shown in Figure 1. The result of elemental distribution analysis (Digital Mapping) of a wavelength-dispersive X-ray spectrometer (WDS) of an electron probe microanalyzer (ΕΡΜΑ) of an alloy. Fig. 2 is an RTB system shown in Fig. 1. The electron beam image of the alloy, the R-rich phase is shown in white, and the R2T14B phase (main phase) is shown in gray. Figure 3 shows the distribution of Fe corresponding to the region of Figure 2. From Figure 2 and Figure 3, we can see that the rich The R phase is smaller than the main phase, and the F-rich phase is less Fe. Figure 4 shows the Nd2 distribution map corresponding to the second graph region. From Fig. 2 and Fig. 4, it can be seen that the R-rich phase is compared with the main phase. The R-rich phase has more Nd, and the fifth graph shows the distribution of Dy corresponding to the second graph region. It can be seen from Fig. 2 and Fig. 5 that the R-rich phase is less than the D-rich R phase in the main phase. number 6 The figure shows the distribution map corresponding to 〇3 in the area of Fig. 2. It can be seen from Fig. 2 and Fig. 6 that the R-rich phase is more than the main phase compared with the 'rich R-rich phase -11 - 200903532. Figure 10 shows the results of elemental distribution analysis by the field emission type electronic micro-survey (FE-EPMA) (Digital drawing ' DigitaI mapping'.) Figure 7 shows the electrons of the RTB-based alloy shown in Figure 1. The beam image 'R-rich phase is shown in white, and the R2T14B phase (main phase) is shown in gray. Figure 8 shows the distribution of Dy corresponding to the region of Figure 7. From Figure 7 and Figure 8, the R-rich phase and the main In contrast, the Dy concentrated region of Dy is formed in the vicinity of the R-rich phase. Furthermore, it can be seen from Fig. 8 that the concentration of Dy is less than that of the Dy concentrated region, and the R-rich phase is less than the main phase. Fig. 9 is a view showing the distribution of Fe corresponding to the region of Fig. 7. It can be seen from Fig. 7 and Fig. 9 that the R-rich phase is less than the main phase, and the R-rich phase has less Fe. The tenth chart does not correspond to the seventh. The distribution map of Nd in the graph area. It can be seen from Fig. 7 and Fig. 10 that the R-rich phase is compared with the main phase, and the R-rich phase is rich in ^^(1). (Manufacturing method) Fig. 1 shows the present invention. RTB For example, it can be cast by the SC method using the alloy manufacturing apparatus shown in Fig. 11. "Alloy manufacturing apparatus" Fig. 1 is a front view showing the overall structure of the alloy manufacturing apparatus of the present embodiment. The alloy manufacturing apparatus 丨 shown in Fig. 11 (hereinafter, the manufacturing apparatus is shown in Table-12-200903532) is roughly composed of a casting apparatus 2, a crushing apparatus 21, and a heating apparatus 3. The heating apparatus 3 is roughly composed of a heater 31 and a container 5. Composition. The container 5 is roughly constituted by a storage container 4 and a switch stage group 32 disposed above the storage container 4. The manufacturing apparatus 1 which is not shown in Fig. 11 is provided with a chamber 6. The tank chamber 6 is composed of a casting chamber 6a and a heat insulating and storage chamber 6b which is disposed below the casting chamber 6a and communicates with the casting chamber 6a. The casting device 2' is housed in the casting chamber 6a, and the heating device 3 is housed in the heat storage/storage chamber 6b. Further, a gate 6e is provided in the heat retention/storage chamber 6b, and the container 5 is transported to the outside of the heat preservation/storage chamber 6b. The heat retention/storage chamber 6b is sealed by the shutter 6e. Further, the crushing device 21 is provided in the casting device 2, and a groove 7 for guiding the cast alloy flakes to the switch type platform group 32 is provided between the casting device 2 and the switch type platform group 3 2 . "Casting device" Fig. 12 is a front view showing the casting device 2 equipped in the manufacturing apparatus 1. The casting apparatus 2' shown in Fig. 2 is substantially a cooling drum 2 2 ' for casting the alloy μ by rapidly cooling the alloy melt L by a water-cooling mechanism (not shown), and supplying the molten alloy L to the cooling drum 2 2 The casting distributor (tun di sh) 23 and the crushing device 21 for crushing the casting alloy μ to form a cast alloy sheet. The crushing device 21 is constituted by, for example, a pair of crushing drums 2 1 a as shown in Fig. 12'. -13- 200903532 "Heating device" Fig. 13 is a front view showing the heating device 3 installed in the manufacturing apparatus 1, and Fig. 14 is a side pattern view. Fig. 15 is a plan view, such as Figs. 13 to 15. The heater 31 constituting the heating device 3 is composed of a heater cover 31a and a heater body 31b attached to the heater cover 31a. The heater cover 31a is designed to radiate heat generated by the heater body 31b to the side surface of the container 5 and prevent heat generated by the heater body 31b from being radiated to the casting chamber 6a. Further, by providing the heater cover 31a, it is possible to prevent the heater body 31b from being damaged when one part of the alloy melt or the cast alloy is dropped from the casting apparatus 2. Further, an opening 3 1 c is provided in the heater 31, and the opening 3p of the bucket 7 is disposed in the opening 3 1 c. Thereby, the cast alloy flakes N dropped from the casting device 2 through the sump 7 can be supplied to the switch-type stage group 3 2 of the container 5 below the heater 31. Further, as shown in Figs. 11 and 13, the heater 31 is disposed along the longitudinal direction of the conveyor belt 51 provided in the heat retention/storage chamber 6b (the moving direction of the container 5). With this configuration, even when the container 5 is moved inside the heat retention/storage chamber 6b, the cast alloy sheet N loaded on the switch type platform group 32 of the container 5 is uniformly insulated. Next, the switch type platform group 3 2 constituting the heating device 3 is integrated with the storage container 4 to constitute the container 5. That is, the -5-200903532 container 5 shown in Figs. 13 to 15 is composed of a storage container 4 and a platform group 32 provided above the storage container 4. In the switching platform group 3 2 ' along the valley: 5, the moving direction sequence is a plurality of switching platforms 33. Further, the member 52' is prevented from being scattered around the switch type plate group 32 by the guide member 52 to prevent the alloy flakes N from being scattered in the heat retention/storage chamber 6b. Each of the switch type platforms 3 3 mounts the sheet S supplied by the prayer device 2 to the holding time set by the heater 31. The cast alloy flakes N are dropped to the storage container 4 after the lapse of time. In each of the stages 33, a plurality of platform plates 33a and a switch platform plate 3 are provided: a closing structure 3 3 b. Each of the switch structures 3 3 b is composed of a rotary shaft 3 3 b mounted on the platform side, and a rotary drive shaft 3 3 b 1 . The inclination angle of each platform plate 3 3 a can control the rotation of the rotary shaft 3 3 b i by each. The degree of each platform plate 3 3 a is at 〇° (the platform plate 3 3 a is horizontal (the state indicated by the broken line in Fig. 3)) to the clockwise direction is about 90° (the platform plate is vertical) 13) The state indicated by the solid line))) is arbitrarily set. Moreover, the switch platform 33 is operated by the switch structure 33b, and the cast alloy sheet N is placed on the platform plate before the predetermined holding time to increase the inclination angle of the platform plate 3 3 a to make the cast alloy thin into the storage container 4 . In addition, the switch platform 3 3 functions as a storage capacity; the cover is such that the heat of the heater 3 1 does not reach the switch of the storage container 4 and the casting alloy is thinly insulated. 33a The driving source tilt angle shown is in the range of two points 33a to pass: 33a, and the μ of the sheet N 4 can prevent -15-200903532 from rising in the temperature inside the storage container 4. Further, a plurality of cooling plates 4a are placed inside the storage container 4. Further, as shown in Figs. 13 and 14, the container 5 is mounted above the conveyor belt 51 (movable device). The container 5 is movable in the left-right direction in Fig. 13 by the conveyor belt 51. "Alloy Casting" Each of Figs. 16 to 19 is a front view showing the operation of the alloy manufacturing apparatus. First, as shown in Fig. 16, the switch platform 33A at the left end of the figure in the switch type platform group 3 2 is placed immediately below the discharge port 7a of the bucket 7 to move the container 5. And 'all the switch display platforms 33 are set to the "off" state. Next, the casting apparatus of Fig. 12 is operated to produce a cast alloy sheet N. First, the alloy melt L is prepared in a melting device (not shown). The temperature of the alloy melt L is adjusted to 130 (TC to 1500 °C depending on the alloy composition. The obtained alloy melt L is conveyed to the casting device 2 together with the refractory crucible 24' from the refractory crucible 24 to the casting The distributor 23 is then supplied from the casting distributor 2 3 to the cooling drum 22, and solidified on the cooling drum 2 2 to form a cast alloy crucible. Thereafter, the casting alloy is kneaded on the opposite side of the casting distributor 23 from the cooling drum 22 Detach the 'casting alloy to the two rotating crushing drums 2 1 a to break it' to make the casting alloy Ν 〇 to the cooling drum 2 2 the average melt supply speed is 1 - 1 cm wide - 200903532 More than 1 每秒g per second, preferably more than 20g per second per 1cm wide, more preferably 25g per second or more per liter, and more preferably less than 100g per second. When the supply speed of the alloy melt L is less than 1 每秒g per second, the alloy melt L cannot be wetted and cooled extensively due to the viscosity of the alloy melt L itself or the wettability with the surface of the cooling drum 22. The drum 22 shrinks on the drum 22, causing the quality of the alloy to change. Further, when the average melt supply speed to the cooling drum 22 is more than 1 000 g per 1 cm wide, the cooling on the cooling drum 22 may be insufficient to cause coarsening of the structure and precipitation of a -Fe. The average cooling rate of the alloy melt on the drum 22 is preferably from 100 to 2 000 ° C per second. If the temperature is above 100 ° C per second, the cooling rate is sufficient, and the precipitation of a -Fe and the coarse R-rich structure are prevented. In addition, if it is 200 or less per second (TC or less, the degree of subcooling does not become excessive, and the cast alloy flakes can be supplied to the heating device 3 at an appropriate temperature. Further, since the cast alloy flakes are not too cold, There is no need to reheat. The average cooling rate can be determined by the difference between the temperature at which the melt immediately contacts the cooling drum and the temperature at which the cooling drum is removed, and the time taken to contact the cooling drum. Further, when the cooling drum 22 is removed, the alloy is cast. The average temperature will subtly vary depending on the subtle difference in the degree of contact between the alloy crucible and the cooling drum 22, the thickness of the cast alloy crucible, etc. The average temperature of the cast alloy crucible is removed from the cooling drum, for example, The radiation thermometer is measured by scanning the surface of the alloy in the width direction from the start to the end of the casting, and is obtained by averaging the obtained measured enthalpy. The average temperature of the cast alloy bismuth from the cooling drum 22 is higher than that of the alloy -17-200903532 The solidification temperature of the r2ti4b phase equilibrium is preferably 100 to 500 ° C, preferably 100 to 400 t. The melting temperature of the R 2 TI 4 B phase ' Nd_Fe · B ternary system is set at 1 150 ° C, but according to N d varies depending on the substitution of the rare earth element, the substitution of the transition element other than Fe, and the type or addition amount of the other added element. The temperature difference between the average temperature of the casting alloy M which is separated from the cooling drum 22 and the solidification temperature in the equilibrium state of the R2T and 4B phases of the cast alloy crucible is insufficient if the loot: On the other hand, when the temperature difference exceeds 500 ° C, the melt is too cold due to the excessive cooling rate. Further, the average temperature □ of the cast alloy 脱离 from the cooling drum 22 fluctuates even in the same casting step (tap), but if the variation is large, the structure and quality are changed. Therefore, the temperature variation within the tapping iron is preferably less than 200 ° C, preferably less than 1 ° C, more preferably 50 ° C, and even more preferably 20 ° C. The average thickness of the cast alloy flakes N is preferably 0.1 mm or more and 1 mm or less. When the average thickness of the sheet is thinner than 0.1 mm, the solidification rate is excessively increased, and the dispersion of the R-rich phase becomes too fine. Further, when the average thickness of the sheet is thicker than 1 mm, the dispersibility of the R-rich phase is lowered due to a decrease in the solidification rate, and precipitation of a - F e is obtained. Next, as shown in Fig. 16, the cast alloy flakes N are sent to the heating device 3 through the hopper 7, and are stacked (placed) on the switch type platform 33A located directly below the discharge port 7a of the hopper 7. At this time, the heater 31 is in the energized state. 'The cast alloy flakes N are placed on the switch platform 3 3 A' and then warmed or warmed by the heater 31. For the switch type platform 33A, the deposited amount of the cast alloy flakes -18-200903532 is preferably set as appropriate according to the area of the deck plate 3 3 a, but the cast flake N is continuously supplied from the casting device 2, although according to the supply speed In either case, the cast alloy flakes N will be replaced by a switch platform 3 3 A. Therefore, when the N amount of the cast alloy flakes corresponding to the switch type platform 33A reaches the set value, 'the container 5 is moved to the left in the figure as shown in Fig. 17', so that the switch type platform 3 3 located on the right side of the switch type platform 3 3 A B. The groove 7 is directly below the discharge port 7a, so that the cast alloy sheet n stacks the switch platform 33B. Thereafter, the container 5 is moved in the same manner as the casting of the casting sheet N, and the casting alloy sheets are sequentially laminated on the respective switching platforms 33C to 33E. The cast alloy thinner deposited on each of the switch platforms 33A to 33E is separately heated or heated by the heater 31. The holding temperature is preferably lower than the temperature of the drum. Specifically, (the drum is separated from the temperature - 1 〇 〇. <:) The range below the drum detachment temperature is preferably (the drum detachment temperature -) is preferably in the range of not less than the drum detachment temperature, and more preferably in the range of °C or more and 900 °C or less. If the holding temperature is 6 〇〇. 〇 ' can fully enhance the coercive force of the R-T-B alloy. Further, if the temperature is kept below 90 °C, the deposition of a-F e and the coarsening of the R-rich group can be prevented. Further, if the temperature of the drum is lowered for any reason, the temperature of the holding temperature is raised to a higher temperature than the temperature at which the drum is removed. Preferably, the amplitude is preferably within 100 ° C and less preferably within 50 ° c. If the temperature rises, the production efficiency will decrease. Moreover, even if it is kept at 10 〇ot, it has the effect of lifting. However, the coarsening of the structure, the fine-grained particle size distribution alloy, and then the direction of the overflow pile moved to the gold thin N pile piece N out of the above 50 ° C 600 temperature woven can heat up the magnetic and micro-19 - 200903532 The flowability of fine powders, even the sintering temperature, will change. Therefore, the effect on subsequent steps must be considered when holding at 1 〇 〇 ° °C. Furthermore, the holding time is preferably 30 seconds or more, 30 seconds to several hours is better, and 30 seconds to 30 minutes is best. If the holding time is more than 30 seconds, the holding force can be sufficiently increased, and the holding time can be kept for several hours, but in terms of production efficiency, it is preferably 30 minutes or less. Next, as shown in Fig. 18, with respect to the remaining switch-type platforms 33F to 33J, the container 5 can be moved by the preparation of the cast alloy flakes N, and the cast alloy flakes N are sequentially stacked on each of the switch platforms. 33F to 33J. When the casting alloy sheets N stacked on the switch platforms 33A to 33D are each subjected to a predetermined holding time, as shown in FIG. 18, the switching platforms 33A are sequentially turned "on" to form a cast alloy sheet. N sequentially falls to the storage container 4. By dropping the cast alloy flakes N into the storage container 4, the heat of the heater 31 does not reach the cast alloy flakes N, thereby terminating the heat retention treatment. As illustrated in Fig. 17, since the cast alloy flakes N are sequentially placed on the respective switch type stages 33A, the holding start time of the cast alloy flakes N on the respective switch type stages 33A... is in each of the switch types. Platform 3 3 has a time difference each. Therefore, in order to maintain the holding time of the cast alloy flakes N on the respective switch type stages 33A, the switch patterns 33A are sequentially "opened", and the cast alloy flakes N are sequentially dropped to the storage container 4. good. The cast alloy flakes N falling on the storage container 4 are sucked by the cooling plate 4a by contact with the cooling plate 4a, thereby cooling the cast alloy flakes N. -20- 200903532 Figs. 19 and 20 show that all of the switch type platforms 33A are in the "open" state, and the cast alloy sheets N are stored in the storage container 4. Then, when the casting and the crushing step are continued by the casting device 2, all the switch patterns 33A are brought into a "closed" state, and the container 5 is moved in the right direction in the drawing, and the casting alloy sheet N is blended. The cast alloy flakes N are sequentially placed on each of the switch platforms 33 A.... When the casting and crushing steps of the casting apparatus 2 are terminated, the entire switching type 33 A is brought into a "closed" state, so that the heat of the heater 31 does not reach the storage container 4. Then, the shutter 6e of the heat storage/storage chamber 6b is opened and the container 5 is transported to the outside of the chamber 6, and the cast alloy flakes N are taken out to terminate the manufacture of the cast alloy flakes N. "Cooling Rate" Next, the cooling rate in the case of manufacturing an R-T-B alloy will be described. In the present invention, the cooling rate is controlled as follows, from the freezing point of the main phase (near 1 1 70 °C) at the temperature just after solidification to 600 °C lower than the freezing point of the R-rich phase. Cooling rate. The cooling rate of the R_T-B alloy is between 100 and 300 ° C per second between 1 000 ° C and 85 ° C. l 〇〇〇 ° C to 850 ° C cooling rate, if faster than the above range, it is considered that Dy will not fully diffuse to the main phase, if it is slower than the above range, then Dy will be excessively diffused It is impossible to form a concentrated portion in the main phase. Further, the cooling rate of the RTB-based alloy is preferably from 300 ° C to 2000 ° C per second from the freezing point of the main phase to 1 ° C. Since the cooling rate of the primary phase freezing point -21 - 200903532 to 1000 °c is in the above range, an R-T-B alloy having a Dy concentration region can be obtained, and high productivity can be obtained. Moreover, the cooling rate of the R-T-B alloy was 850. (: It is preferable to temporarily become 100 ° C or less in the region of 600 ° C. Since the cooling rate of 850 ° C to 600 C is temporarily set to the above range, Dy contained in the r-rich phase can be sufficiently diffused to Since the main phase is adjacent to each other, the RTB-based alloy having the Dy Φ shrinkage region and the Fe-force is more easily produced. The RTB-based alloy and the R_T_B-based alloy flake of the present embodiment are formed in the vicinity of the R-rich phase and have Compared with the composition ratio Dy concentrated Dy concentrated region, it is a rare earth permanent magnet which has high coercive force and excellent magnetic properties. That is, the RTB alloy of the present embodiment has, for example, a comparison with Fig. 21 Figure 24 shows a higher coercive force of an RTB-based alloy without a Dy-concentrated area. Figures 21 to 24 show FE-EPMA (Field Emission-type Electron Micro) as an example of an RTB-based alloy without a Dy-concentrated region. The result of the element distribution analysis. The R_T_B alloy shown in Fig. 21 to Fig. 24 was obtained by the SC method. The composition of the RTB alloy was mass ratio: N d 2 3 %, D y 9 %, B 1 %, C ο 1 %, ga 0 · 2 %, the rest is F e ° Figure 21 is an electron beam image of the r_T_b alloy without Dy concentrated region. The R-rich phase is shown in white, the r2t14b phase (main phase) is shown in gray. Figure 22 shows the Dy corresponding to the 21st region. The distribution map. From -22-200903532 2 1 and 2 2, it is known that there is no D y concentrated D y concentrated region compared with the main phase, and the Dy concentration is compared with the main phase, and the R-rich phase is also Fig. 23 is a view showing the distribution of Fe corresponding to the region of Fig. 21. It can be seen from Fig. 21 and Fig. 23 that the Fe of the R-rich phase is smaller than that of the main phase, and the figure 24 corresponds to the region of the 21st map. The distribution map of Nd. It can be seen from Fig. 21 and Fig. 24 that the R-rich phase has more Nd than the main phase. (Preparation of RTB rare earth permanent magnet) The R - T - B rare earth of the present invention is produced. In the case of the RTB-based alloy of the present invention, the fine powder of the RTB-based rare earth permanent magnet is used. The fine powder for the RTB-based rare earth permanent magnet of the present invention is, for example, a sheet made of the RTB-based alloy of the present invention. Method for finely pulverizing hydrogen absorption after being decomposed by hydrogen using a pulverizer such as a jet honing machine The hydrogen disintegration step is preferably carried out, for example, by a hydrogen absorption step of maintaining a hydrogen atmosphere maintained at a predetermined pressure in advance. Next, the obtained fine powder of the RTB-based rare earth permanent magnet is molded by using a transverse magnetic field, for example. The RT-based rare earth permanent magnet is obtained by press molding and sintering in a vacuum. The R-T_B rare earth permanent magnet fine powder and the RTB rare earth permanent magnet of the present embodiment are the R-T_B system of the present invention. If the alloy is made, it is high in magnetic strength and excellent in magnetic properties. [Embodiment] [Example 1] -23- 200903532 The mass ratio is: N d 2 3 %, D y 9 %, B 0 · 9 8 %, C ο 1 %, G a 0 · 2 %, and The remainder was weighed with the raw materials of F e, and the alumina melt was prepared by melting in a high-frequency melting furnace in an argon atmosphere at a pressure of 1 atmosphere. Next, the casting of the alloy melt to the manufacturing apparatus shown in Fig. 1 is cast by the SC method. The peripheral speed of the cooling drum during casting was 1.3 m/s, and the average melt supply rate to the cooling drum was 30 g per 1 cm width per second. The average temperature of the cast alloy block leaving the cooling drum was 850 C. The cooling rate of this alloy is: from the freezing point of the main phase to 700 ° C / sec for 1 〇〇〇 ° C, 200 ° C / sec from 1 000 ° C to 850 ° C, and 85 ° C to 7 80 ° c is 5 (TC / sec. Thereafter, it is maintained at a temperature of about 78 ° C for 300 seconds on the switch platform using the manufacturing apparatus of Fig. 11, and then cooled to 6 冷却 at a cooling rate of 0.1 ° C / sec. The r-Τ-lanthanide alloy flakes of Example 1 were produced at a temperature below ° C. The average thickness of the alloy at this time was 0.3 mm. [Example 2] An alloy melt was prepared using the same raw materials and equipment as in Example 1. Using the same casting apparatus as in Example 1, the obtained alloy melt was set to: the circumferential speed of the cooling drum at the time of casting 〇 _ 8 7 m / sec, and the average melt supply speed supplied to the cooling drum per centimeter The casting is performed at a width of 3 〇g per second and the casting alloy block is separated from the cooling drum at an average temperature of 8 8 〇 ° C. The cooling rate of the alloy is from the freezing point of the main phase to 1 0 0 〇 ° C for 7 00. °C / sec, 1 000 °C to 850 °c is 2 0 0 °C / sec, 8 5 0 °C to 7 80 °c is -24- 200903532 10C / sec, thereafter Using the manufacturing apparatus of the table 11, the RTB-based alloy flakes of Example 2 were produced without using a switch platform and cooled to a temperature of 600 ° C or less at a cooling rate of 0-1 ° C / sec. The average thickness of the alloy at this time was 〇 45 mm. The R_T_B-based alloy flakes of Example 1 and Example 2 were obtained by elemental distribution analysis (Digital Mapping) (Wave Surface Analysis) of WDS (wavelength dispersion X-ray spectrometer). EpMA and FE-EPMA. The RTB-based alloy flakes of Example 1 and Example 2 were each formed in the vicinity of the R-rich phase to form an R-rich phase and a Dy concentrated phase which was concentrated with the main phase. Further, Example 1 And the RTB-based alloy flakes of Example 2, the Dy concentration of either of them was less than the Dy concentrated region, and the R-rich phase was less than the main phase. (Comparative Example 1) The same raw materials as in Example 1 were used. The apparatus was used to prepare an alloy melt. Using the same casting apparatus as in Example 1, the obtained alloy melt was set to: the peripheral speed of the cooling drum at the time of casting was 0.65 m/s, and the average melt supply speed of the cooling drum was 1 cm. 1 5 g/sec wide, while the cast alloy block is out of the cold While the average temperature of the drum was 700 ° C, the RTB-based alloy flakes of Comparative Example 1 were obtained. The cooling rate of the alloy was: from the freezing point of the main phase to 1 〇〇〇 ° C of 700 ° C / Seconds, l°°C to 700°C is 400°C/sec, 7〇〇°C to 600°C is j 〇r / sec, after which, the manufacturing device of Figure 11 is used instead of the switch platform Cool to below 600 °C at a cooling rate of 0.1 °C per second. At this time, the average thickness of the alloy is 〇.30 mm. -25- 200903532 About the obtained R-Τ-lanthanide alloy flakes of Comparative Example 1 WDS_EPMA and FE-EPMA were subjected to elemental distribution analysis (digital side; Digital mapping) (surface analysis). As a result, Comparative Example 1 [B-based alloy flakes did not form a concentrated region in which Dy was concentrated by Dy compared with the composition ratio. Salt is considered to be one of the reasons why the temperature of the cast alloy block is out of the cooling drum in Comparative Example 1 and the alloy is excessively rapidly cooled in the cooling roll, and the cooling rate of I OOOt to 70 0 t is too fast, and Dy and Nd cannot be sufficiently diffused. It is impossible to form a concentration gradient. Next, as shown below, magnets were produced using the obtained R-Τ-lanthanum alloy sheets of Example 1 and Examples and Comparative Example 1. First, the R-T alloy flakes of Example 1 and Example 2 and Comparative Example 1 were subjected to hydrogenolysis. Hydrogenolysis is carried out by absorbing hydrogen in hydrogen at 2 atmospheres after each R-Τ-Β-separated sheet, and then heating to 500 ° C residual hydrogen in a vacuum, followed by adding 0.07 mass% of zinc stearate and using nitrogen gas. The honing mill is carried out by means of fine pulverization. The fine particle size obtained by fine pulverization was not about 5.0 m as measured by laser diffraction. Next, the obtained powder was formed by press molding at a pressure of 〇8 t/cm 2 in a nitrogen atmosphere using a molding machine in a transverse magnetic field, and then the formed body was heated from the chamber at a temperature of 1.33 x 10 5 hPa at 500 °. C and 8 00 ° C were each kept for one hour to remove residual hydrogen from stearic acid. Thereafter, the temperature was raised to 1 〇 30 ° C of the sticking temperature, and 3 was maintained to obtain a burnt body. Thereafter, the obtained burned bodies were heat-treated at 80 ° C and 530 ° C for one hour in an argon atmosphere, and each of the examples 1 to 2 was obtained, and the comparative examples were obtained. magnet. In the drawing, the R-shrinkage is caused by the cylinder of the sample 2 - B system to extract the 100% of the body. The magnetic properties of the magnets of Example 1 and Example 2, which were obtained by measuring B H C u r v e T r a c e r , were measured by B C Η hysteresis curve meter (B H C u r v e T r a c e r ). The results are shown in Table 1 and Figure 25. Fig. 25 is a graph showing the coercive force (η Cj ) of the magnets of Example 1, Example 2 and Comparative Example 1, and the vertical axis shows the coercive force 'horizontal axis' level. Further, in Fig. 25, the symbol 〇 indicates the coercive force of the first embodiment and the second embodiment, and the symbol ▲ indicates the coercive force of the comparative example 1. -27- 200903532 [Table 1]
Hcj (kOe) Hk/Hcj (%) BHmax (MGOe) Br (kG) 實施例1 33.34 56.40% 32.76 11.54 34.19 54.87% 32.61 11.54 33.69 56.39% 33.20 11.62 33.99 56.73% 33.36 11.63 34.10 56.43% 33.27 11.63 33.35 55.58% 32.41 11.47 33.18 57.28% 33.27 11.61 33.44 57.95% 33.72 11.71 33.67 56.62% 33.08 11.58 33.13 57.53% 33.22 11.61 實施例2 32.61 58.71% 33.26 11.64 33.07 58.03% 32.96 11.58 33.31 58.49% 33.66 11.68 33.74 57.79% 33.15 11.61 33.01 58.33% 33.03 11.58 33.31 58.09% 32.82 11.60 33.31 58.84% 32.87 11.57 33.69 57.53% 33.00 11.60 33.47 59.14% 33.35 11.65 33.04 60.06% 33.22 11.66 比較例1 31.44 52.30% 32.22 11.45 31.73 52.22% 32.73 11.54 31.28 53.54% 32.73 11.53 31.38 53.08% 33.50 11.70 31.42 52.93% 32.92 11.57 另外，表1中之「（ BH) max」爲最大能源面積，「 Br」爲殘留磁束密度，「Hcj」爲保磁力，「Hk/Hcj」磁 滯之方形性。 如表2及第2 5圖所示，可確認實施例1至實施例2 -28- 200903532 與由未形成Dy濃縮區域之R-Τ-Β系合金所製得之比較例 1相較，具有高保磁力「Hcj」。此種保磁力之差起因於合 金狀態之濃度分佈於粉碎及燒結成爲磁鐵後亦賦予影響。 咸認原因之一爲本發明之合金由於存在Dy濃縮區域，而 此濃縮區域亦殘留於磁鐵之結晶粒子內，因而殘留富R而 不能有效提升保磁力之Dy少。 【圖式簡單說明】 第1圖係顯示本發明之R-T_B系合金之一例的照片， 即藉由掃描型電子顯微鏡（SEM )觀察R-T-B系合金的薄 片之剖面時的照片。 第2圖係第1圖所示之R-T-B系合金之電子束圖像。 第3圖表示對應第2圖區域之Fe之分佈圖。 第4圖表示對應第2圖區域之Nd之分佈圖。 第5圖表示對應第2圖區域之Dy之分佈圖。 第6圖表示對應第2圖區域之Ga之分佈圖。 第7圖爲第1圖所示之R-T_B系合金之電子束圖像。 第8圖表示對應第7圖區域之Dy之分佈圖。 第9圖表示對應第7圖區域之Fe之分佈圖。 第10圖表示對應第7圖區域之Nd之分佈圖。 第11圖表示本發明之實施方式的合金製造裝置之構 成之正面模式圖。 第12圖表示裝備於合金製造裝置之鑄造裝置之正面 模式圖。 -29- 200903532 第]3圖表示裝備於合金製造裝置之加熱裝置之正面 模式圖。 第14圖表示裝備於合金製造裝置之加熱裝置之側面 模式圖。 第15圖表示裝備於合金製造裝置之開關式平台（ stage)與儲藏容器（container)之平面模式圖。 第16圖爲說明合金製造裝置運作之正面模式圖。 第17圖爲說明合金製造裝置運作之正面模式圖。 第18圖爲說明合金製造裝置運作之正面模式圖。 第19圖爲說明合金製造裝置運作之正面模式圖。 第20圖爲說明合金製造裝置運作之側面模式圖β 第21圖爲不具有Dy濃縮區域之R-T-B系合金之電子 束圖像。 第22圖表示對應第21圖區域之Dy分佈圖。 第23圖表示對應第21圖區域之Fe分佈圖。 第24圖表示對應第21圖區域之Nd分佈圖。 第2 5圖表示實施例1、實施例2及比較例1之磁鐵的 保磁力（Hcj )之圖。 t主要元件符號說明】 1:製造裝置（合金製造裝置） 2 :鑄造裝置 3 :加熱裝置 4 :儲藏容器 -30- 200903532 4a :冷卻板 5 :容器 6:箱室(chamber) Ί ··斗槽(hopper ) 7 a :斗槽之排出口 2 1 :破碎裝置 3 1 :加熱器 3 1 c :開口部 3 3 ::開關式平台 3 3 a :平台板 3 3 b :開關機構 51:傳送帶（可動裝置） L :合金熔液 N :鑄造合金薄片 6a :鑄造室 6b :保溫·儲藏室 6 e :閘門（g a t e ) 2 1 a :破碎滾筒 23 :澆鑄分配器 24 :坩堝 3 1 a :加熱器蓋 3 1 b :加熱器本體 3 2 :平台組 3 3 b :開關結構 -31 200903532 3 3 b ,:迴轉軸 5 2 :導向構件Hcj (kOe) Hk/Hcj (%) BHmax (MGOe) Br (kG) Example 1 33.34 56.40% 32.76 11.54 34.19 54.87% 32.61 11.54 33.69 56.39% 33.20 11.62 33.99 56.73% 33.36 11.63 34.10 56.43% 33.27 11.63 33.35 55.58% 32.41 11.47 33.18 57.28% 33.27 11.61 33.44 57.95% 33.72 11.71 33.67 56.62% 33.08 11.58 33.13 57.53% 33.22 11.61 Example 2 32.61 58.71% 33.26 11.64 33.07 58.03% 32.96 11.58 33.31 58.49% 33.66 11.68 33.74 57.79% 33.15 11.61 33.01 58.33% 33.03 11.58 33.31 58.09% 32.82 11.60 33.31 58.84% 32.87 11.57 33.69 57.53% 33.00 11.60 33.47 59.14% 33.35 11.65 33.04 60.06% 33.22 11.66 Comparative Example 1 31.44 52.30% 32.22 11.45 31.73 52.22% 32.73 11.54 31.28 53.54% 32.73 11.53 31.38 53.08% 33.50 11.70 31.42 52.93% 32.92 11.57 In addition, "(BH) max" in Table 1 is the maximum energy area, "Br" is the residual magnetic flux density, "Hcj" is the coercive force, and "Hk/Hcj" is the squareness of the hysteresis. As shown in Table 2 and Figure 25, it can be confirmed that Example 1 to Example 2-28-200903532 are compared with Comparative Example 1 obtained from an R-Τ-lanthanum alloy in which no Dy concentration region is formed. High magnetic force "Hcj". This difference in coercive force is caused by the concentration distribution in the alloy state, which is also affected by the pulverization and sintering. One of the reasons for this is that the alloy of the present invention has a Dy-concentrated region, and this concentrated region remains in the crystal particles of the magnet, so that the residual R is rich and the Dy which does not effectively increase the coercive force is small. BRIEF DESCRIPTION OF THE DRAWINGS Fig. 1 is a photograph showing an example of an R-T_B-based alloy of the present invention, that is, a photograph of a cross section of a sheet of an R-T-B-based alloy observed by a scanning electron microscope (SEM). Fig. 2 is an electron beam image of the R-T-B alloy shown in Fig. 1. Fig. 3 is a view showing the distribution of Fe corresponding to the area of Fig. 2. Fig. 4 is a view showing the distribution of Nd corresponding to the area of Fig. 2. Fig. 5 is a view showing the distribution of Dy corresponding to the area of Fig. 2. Fig. 6 is a view showing the distribution of Ga corresponding to the area of Fig. 2. Fig. 7 is an electron beam image of the R-T_B alloy shown in Fig. 1. Fig. 8 is a view showing the distribution of Dy corresponding to the area of Fig. 7. Fig. 9 is a view showing the distribution of Fe corresponding to the area of Fig. 7. Fig. 10 is a view showing the distribution of Nd corresponding to the area of Fig. 7. Fig. 11 is a front view showing the configuration of an alloy manufacturing apparatus according to an embodiment of the present invention. Fig. 12 is a front view showing the casting apparatus equipped in the alloy manufacturing apparatus. -29- 200903532 Fig. 3 is a front view showing the heating device equipped in the alloy manufacturing apparatus. Fig. 14 is a side view showing the heating device equipped in the alloy manufacturing apparatus. Fig. 15 is a plan view showing a switch stage and a container equipped in an alloy manufacturing apparatus. Figure 16 is a front view showing the operation of the alloy manufacturing apparatus. Figure 17 is a front view showing the operation of the alloy manufacturing apparatus. Figure 18 is a front view showing the operation of the alloy manufacturing apparatus. Figure 19 is a front view showing the operation of the alloy manufacturing apparatus. Fig. 20 is a side view showing the operation of the alloy manufacturing apparatus. Fig. 21 is an electron beam image of an R-T-B alloy having no Dy concentrated region. Fig. 22 shows a Dy distribution map corresponding to the area of Fig. 21. Fig. 23 is a view showing the Fe distribution map corresponding to the area of Fig. 21. Fig. 24 is a view showing the Nd distribution map corresponding to the area of Fig. 21. Fig. 25 is a view showing the coercive force (Hcj) of the magnets of the first embodiment, the second embodiment, and the comparative example 1. t Main component symbol description] 1: Manufacturing device (alloy manufacturing device) 2: Casting device 3: Heating device 4: Storage container -30- 200903532 4a: Cooling plate 5: Container 6: Chamber Ί · · Bucket (hopper ) 7 a : Discharge port 2 1 : Crusher 3 1 : Heater 3 1 c : Opening 3 3 :: Switching platform 3 3 a : Platform plate 3 3 b : Switching mechanism 51: Conveyor belt ( Movable device) L: alloy melt N: cast alloy flake 6a: casting chamber 6b: heat preservation and storage chamber 6e: gate 2 1 a : crushing drum 23: casting distributor 24: 坩埚3 1 a : heater Cover 3 1 b : heater body 3 2 : platform group 3 3 b : switch structure -31 200903532 3 3 b ,: rotary shaft 5 2 : guide member